Document [original]

Inﬂuence of α-Precipitate Orientation and Distribution on

the Deformation Behavior of the Additively Manufactured

Metastable β-Titanium Alloy Ti-5553 Assessed by Cyclic

Nanoindentation

Erika Gabriele Alves Alcântara and Claudia Fleck*

1. Introduction

Titanium (Ti) alloys are in common use for orthopedic implants

because of their excellent combination of mechanical strength,

corrosion resistance, and biocompatibility.

[1,2]

The metastable

β-alloy Ti–5Al–5V–5Mo–3Cr (wt%, Ti-5553) has recently found

growing interest for load-bearing implants. As compared to

the up-to-date gold standard, the (αþβ)-

alloy TiAl6V4, Ti-5553, exhibits a better

combination of ductility, toughness, and

Young’s modulus.

[3]

Due to daily activities, load-bearing

implants and, speciﬁcally, joint replace-

ments have to sustain very high loads over

millions of cycles.

[4]

Besides biocompatibil-

ity, implant materials hence need to exhibit

sufﬁcient fatigue resistance. It is, therefore,

of prime importance to understand their

cyclic deformation and fatigue behavior.

Previous studies have reported on the

inﬂuence of processing variables and

follow-up treatments on the microstructure

and fatigue behavior of additively

[5]

and

conventionally

[6–14]

manufactured Ti-5553.

Only a few of these have also investigated

the cyclic deformation mechanisms.

[12,14]

Results are ambiguous, due to different load-

ing conditions and different microstructures

of the tested materials. For example, during

pure compression fatigue loading, Ti-5553

with a β-annealed microstructure exhibited

initial cyclic softening due to dislocation annihilation, which

reduces constraints on dislocation mobility, and evolution of twin

structures, such as detwinning and twin boundary degradation.

With further loading, the material entered a saturation state, attrib-

uted to the ﬂip-ﬂop movement of dislocation dipoles.

[13]

Under fully

reversed tension–compression loading, Ti-5553 with an (αþβ)-

microstructure showed plastic strain incompatibility between the

αand βphases. In strain-controlled loading, the cyclic deformation

behavior further depends on the strain amplitude. At low levels,

some studies observed moderate hardening at the beginning of

loadingfollowedbyasaturationstage,

[12]

while others reported only

softening.

[14]

At intermediate amplitudes, moderate hardening pre-

ceded softening,

[12]

and at high amplitudes, pronounced softening

occurred.

[12,14]

Cyclic hardening was attributed to the activation and

interactions of multiple slip systems in the primary α(α

) phase and

the impingement of dislocation movement by α

/βboundaries.

Softening was explained to be the result of massive dislocation

annihilation, mainly of pre-existing dislocations, and cross-slip

of dislocations in the α

-precipitates and the β-matrix.

[12,14]

Conventional fatigue testing of macrospecimens requires a lot

of material and a high number of specimens. Cyclic nanoinden-

tation offers a fast and relatively simple approach to overcome

E. G. Alves Alcântara, C. Fleck

Faculty III Process Sciences

Institute of Materials Science and Technology

Fachgebiet Werkstofftechnik/Chair of Materials Science & Engineering

Technische Universität Berlin

Strasse des 17. Juni 135, 10623 Berlin, Germany

E-mail: claudia.ﬂeck@tu-berlin.de

The ORCID identiﬁcation number(s) for the author(s) of this article

can be found under https://doi.org/10.1002/adem.202301095.

VCH GmbH. This is an open access article under the terms of the Creative

Commons Attribution License, which permits use, distribution and

reproduction in any medium, provided the original work is properly cited.

DOI: 10.1002/adem.202301095

The metastable β-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 biocompati-

bility, implant materials need to exhibit sufﬁcient fatigue resistance. Herein, cyclic

nanoindentation is applied up to a maximum cycle number of 10

to elucidate the

inﬂuence of the local phase distribution on the cyclic deformation behavior of Ti-

5553, made by laser powder bed fusion of metals (LPBF-M), in the (αþβ)-solution

annealed state. By combining the localized cyclic mechanical loading and high-

resolution transmission electron microscopy, the inﬂuence of the presence and

orientation of α

-precipitates within the β-grains on the cyclic deformation behavior

and mechanisms is unraveled. α

-phase orientation and distribution signiﬁcantly

contribute to the effectiveness of the precipitates as barriers to dislocation motion.

A high density and trapping of dislocations are observed at α/βinterfaces. The

occurrence and size of the pile-up surrounding the indents are correlated with

the cyclic deformation behavior, and, thus, with the presence of α

-precipitates. The

gained improved knowledge of the phase-dependent deformation behavior helps to

better understand the fatigue performance of this alloy also on the macroscale.

RESEARCH ARTICLE

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these limitations.

[15]

Furthermore, the method allows probing dif-

ferences in the properties of single phases as well as the inﬂuence

of the local microstructure in a quasi-nondestructive way.

[16–18]

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.

[19–22]

The low num-

ber of publications reporting on bulk materials usually address

the cyclic deformation behavior up to relatively low maximum

numbers of cycles in the range of 10–300 only.

[23–26]

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 vol-

ume fractions were identiﬁed.

[23]

The progression of indentation

depth with the number of cycles exhibited two distinct regions

for solution-treated and aged AZ61 Mg alloy

[24]

and duplex

stainless steel, cyclically nanoindented in load control up to

300 cycles.

[25]

A primary or transient stage where the indentation

depth rate decreased with the number of cycles preceded a sec-

ondary “steady-state”region, suggesting a balance between cyclic

hardening and softening. In displacement-controlled tests up to a

maximum number of cycles of 100, TiAl6V4 with a duplex micro-

structure with more than 90% α-phase exhibited cyclic softening

for both, primary and lamellar α-phase, indicated by a decrease in

peak stress.

[26]

To the best of our knowledge, there are only two

reports so far on cyclic nanoindentation fatigue of metals up to a

much higher maximum cycle number of 10

[27,28]

The tests per-

formed on cross sections of struts extracted from A356.0 alumi-

num alloy open-cell foam

[27]

and of Mg–SiC nanocomposites

[28]

showed signiﬁcant inﬂuences of the phase composition in a

micrometer-sized interaction volume below the indent on the

cyclic deformation behavior.

As reviewed above, 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, we performed cyclic nanoindentation tests on this

alloy in the (αþβ)-solution-annealed state up to a maximum

cycle number of 10

, yielding information on the cyclic deforma-

tion behavior in loading regimes that are relevant to many

applications. By combining cyclic nanoindentation and high-

resolution electron microscopy, we unravel the inﬂuence of dif-

ferent α-phase orientations and distributions within the β-grains

on the deformation mechanisms under cyclic loading. The

gained improved knowledge of the phase-dependent deforma-

tion behavior will help us to better understand the fatigue

performance of this alloy also on the macroscale, ultimately pav-

ing the way to predict the fatigue response in silico for speciﬁc

microstructures.

2. Experimental Section

2.1. Material

Specimens for nanofatigue tests were extracted from the gauge

length of a standard cylindrical fatigue specimen, provided by the

Chair of Machine Tools and Production Engineering of

Technische Universität Berlin (Figure 1). The specimen was pro-

duced by laser powder bed fusion of metals (LPBF-M) with an

SLM Solutions 250 H machine (MTT Technologies GmbH,

Germany) with its longitudinal axis at a 90° angle to the building

platform from Ti-5553 powder made by plasma atomization

(grain size 15–45 μm; D

=34 μm; AP&C Advanced Powders

& Coatings, Quebec, Canada). Following LPBF-M, the specimen

was heat treated to generate a binary (αþβ)-microstructure.

[29]

The chemical compositions of the powder, as given by the sup-

plier, and of the specimen after LPBF-M and after heat treatment,

obtained by energy-dispersive X-ray spectroscopy (EDAX

PV9800, installed on a CamScan REM Serie 2, Obducat,

Sweden) at an accelerating voltage of 20 keV are given in

Table 1 together with the used LPBF-M parameter settings.

2.2. Preloading Microstructural Investigation

Longitudinal and cross sections (Figure 1) were ground on SiC

abrasive paper down to 500 grit and polished with diamond

suspension down to a grain size of 9 μm. Grain size and phase

distribution were evaluated by light microscopy (DMR, Leica,

Germany) after etching with Kroll’s reagent.

[30]

The sections were

then repolished for 10 min with 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 diffrac-

tion (EBSD). The EBSD measurements were performed at the

Central Electron Microscopy Unit (ZELMI) of Technische

Universität Berlin in a ﬁeld-emission scanning electron micro-

scope (SEM) DSM 982 GEMINI (ZEISS, Oberkochen,

Germany) operated at 15 kV with a step size of 100–400 nm.

Image analysis to obtain phase maps and inverse pole ﬁgures

(IPFs) was done with the software OIM Analysis 6.0 (EDAX/

AMETEK, Mahwah, USA).

2.3. Nanofatigue Experiments

Nanofatigue experiments were performed on a Hysitron

Triboindenter TI950 (Bruker Corporation, Massachusetts, USA)

Figure 1. Specimens used for microstructural investigations and nanofa-

tigue tests, prepared from the gauge length of a standard cylindrical

fatigue specimen produced by LPBF-M.

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using a Berkovich tip. The surface of a cross section through an

LPBF-M specimen was prepared in the same way as described

above for the metallographic sections used for the EBSD meas-

urements. 24 indents were placed in a square map, spaced

equally at a distance of 30 μm to avoid mutual interactions.

Four additional indent positions were selected by visual inspec-

tion of etched surfaces; the regions of interest were distinguished

by local variations in α-phase occurence. The identiﬁed positions

were then indented cyclically after repolishing to remove the

roughness induced by etching.

For nanofatigue loading, each site was indented cyclically to a

maximum cycle number of 10

at a frequency, f, of 201 Hz. The

minimum force, P

mín

,was258μNandthemaximumforce,P

max

was 2888 μN, corresponding to a force amplitude, P

, of 1315 μN

and a mean force, P

, of 1573 μN(indents1–24). The four posi-

tions placed in deﬁned regions of interest (addressed in the follow-

ing as indents A, B, C, D) were loaded with P

mín

=456 μNand

max

=2946 μN(P

=1245 μN, P

=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 was 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 (Figure 2a) were inserted at regular intervals to yield

force/indentation depth hysteresis loops with sufﬁcient data

points (Figure 2b). Further, for indents 1–24, static “holding”

segments at 10% of the maximum load were introduced before

and after the measurement cycles for thermal drift correction.

[31]

The force/indentation depth–hysteresis loops were evaluated

according to protocols used in classical fatigue testing by custom-

ized code based on Python.

[32]

The cyclic deformation behavior

was characterized by the development of the plastic indentation

depth amplitude, D

a,p

, determined as the half width of the

hysteresis loop at mean force. The development of the ratios

of D

min

to D

max

over the cycle number, N, and the change of

plastic deformation between cycles, ΔD

min

, induced by repeated

application of the force amplitude over N, and represented by the

change in minimum depth (D

min

) between consecutive measure-

ment cycles, was analyzed to evaluate the cyclic creep behavior.

2.4. Microstructural and Morphological Investigation of

Fatigue-Induced Nanoindent Characteristics

The four nanoindents “A”to “D,”placed in deﬁned regions of

interest, were imaged in a high-resolution SEM (HRSEM;

Gemini SEM500 NanoVP, Zeiss, Oberkochen, Germany) in

Table 1. LPBF-M parameter settings used to manufacture the standard cylindrical fatigue specimen, together with the chemical composition (wt%) of the

Ti-5553 powder, as given by the supplier, and of the specimen after LPBF-M and heat treatment, obtained from energy-dispersive X-ray spectroscopy.

LPBF-M building parameters

Layer thickness D

[μm] Laser power P

[W] Scanning speed v

Hatch pattern Hatch distance Δ

[mm] Focus position

[mm s

1

45 250 1000 Stripes 0.12 0

Condition Element [wt%]

Ti Al V Mo Cr S

Powder Balance 4.86 4.86 4.97 2.8 –

After LPBF-M 9.46 4.39 2.49 2.88 3.66

After heat treatment 7.81 4.49 2.52 2.54 3.44

Figure 2. 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 (D

max

), the minimum depth (D

min

), and the plastic

indentation depth amplitude (D

a,p

) are determined.

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the backscattered electron mode at a voltage of 8 kV at a working

distance of 7.6 mm. Microstructure and dislocation structure and

density in the volume directly beneath these indents were inves-

tigated by transmission electron microscopy (TEM), using a

Tecnai G

20 S-TWIN (FEI Company, OR, USA) at an operating

voltage of 200 kV in the bright-ﬁeld mode. For TEM observation,

thin foils were prepared using the focused ion beam (FIB) tech-

nique (FEI Helios NanoLab 600; Field Electron and Ion

Company, Hillsboro, USA). One long edge of the foil was ori-

ented parallel to the indentation direction, and the section was

placed as precisely as possible through the tip of each nanoind-

ent. To protect the foil against plastic deformation and the sam-

ple surface from the gallium ions, the sample was covered with a

thin platinum layer before ion-beam milling. Although we can-

not be sure that our TEM foils were placed right through the tip,

we are sure that they passed very close to it, given that the maxi-

mum indentation depth after fatigue loading observed from the

triangular proﬁle in the TEM micrographs (compare Figure 11)

matched the range of depths of the four indents “A”to “D”

(182–261 nm). All SEM, FIB, and TEM work was performed

at the Central Electron Microscopy Unit (ZELMI) of Technische

Universität Berlin.

To allow qualitative and quantitative evaluation of indent and

pile-up morphology and size, all cyclic nanoindents and the

surrounding surfaces were imaged using the scanning probe

microscopy (SPM) mode of the nanoindenter. Both topography

and gradient images were evaluated qualitatively and quantitatively.

The projected area (A

), the outer pile-up area (A

p-u

)andthepile-up

volume (V

p-u

) were determined after background subtraction using

Fiji,

[33,34]

where the best ﬁtforA

and A

p-u

was found by visual

inspection of the SPM topography image, as shown for a typical

example in Figure 3.V

p-u

was determined by the integration of

all pixels under the pile-up area. Nanoindent and pile-up proﬁles

were measured using Gwyddion,

[35]

from which the maximum

pile-up height (h

p-u,max

) was obtained, as exempliﬁed in Figure

S1, Supporting Information.

3. Results

3.1. Microstructure

The material has a biphasic microstructure with small

-precipitates highly dispersed within the grains and at the grain

boundaries of the retained β-matrix (Figure 4, and 5e,f ). In low-

magniﬁcation light micrographs (Figure 4a-d), the grain bound-

aries appear as bright stripes, some of which contain black lines.

Higher magniﬁcations (Figure 4e,f) reveal that these bright

stripes consist of β-phase and that the black lines are α

-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 βand α

, respectively, for both transverse

and longitudinal sections (Figure 5a,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 IPF (Figure 5c,d) show no crystal-

lographic texture.

By light microscopy, lighter and darker regions are visible in

both cross and longitudinal sections. Each of these regions com-

prises several grains (Figure 4a,b). EBSD-IPF views (Figure 5e)

and higher-magniﬁed SEM micrographs (Figure 5f ) reveal that

the differences in gray value are due to different densities,

shapes, and amounts of sections of α

-precipitates in the ﬁeld

of view, as shown exemplarily in Figure 6. The different views

reveal that the precipitates are acicular, with their diameters vary-

ing along their length. Hence, the differences in the appearance

of α

come from their different orientation and, thus, the appear-

ance of the exposed sectioned plane. Even though the surface

comprises regions where the α

-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.

3.2. Cyclic Deformation and Creep Behavior

The cyclic deformation behavior is displayed in Figure 7 by plots

of the plastic displacement amplitude D

a,p

versus N. On a mac-

roscopic view, we observe stages of hardening, saturation, and

softening. However, within each of these three stages, ﬂuctua-

tions of D

a,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 D

a,p

This stage is characterized by moderate alterations between hard-

ening and softening for more than 50% of these indents, by soft

alterations for a quarter of the indents and by strong alterations

Figure 3. Morphological evaluation of cyclic nanoindents: a) SPM topography image of a typical example, with the α- and β-phase appearing in dark and

bright gray, respectively; b) typical selection of the projected indent area; c) pile-up area, determined by visual inspection.

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for about a ﬁfth. Indents 3 and 16 (Figure 7b) are examples of

moderate and soft alterations, respectively. Only about 3% of

the nanoindents, for instance indent 1 (Figure 7b), exhibit soft-

ening following pronounced hardening in the ﬁrst cycle and sub-

sequent strong ﬂuctuations. For most nanoindents, the initial

hardening or softening is followed by weak further softening

for loading up to 10

cycles: more than half of the indents

(54%) show an increase in D

a,p

with pronounced ﬂuctuations,

18% with small ﬂuctuations and 7% without ﬂuctuations (com-

pare indents 1 and 16 in Figure 7b). Only about one-ﬁfth of the

nanoindents, for example, indent 3 (Figure 7b), exhibit a low

amount of further hardening: 18% show a decrease in D

a,p

with

ﬂuctuations, and 3% without ﬂuctuations. During further

loading to the maximum number of cycles of 10

, the curves

of most indents (75%) decrease with ﬂuctuations, 22% increase

with ﬂuctuations, and 3% of them present a pronounced increase

with ﬂuctuations.

The progressions of the ratios of D

min

to D

max

and of ΔD

min

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 D

min

max

over N

(Figure 8) progress in similar ways, with an overall increase

in D

min

max

, but they differ in the extent of the increase: some

curves progress to higher and others to lower values.

ΔD

min

represents the incremental, nonreversible deformation,

induced by repeated loading, over the entire loading history. As

Figure 4. a–f) Light micrographs of polished and etched cross (a,c,e) and longitudinal (b,d,f) sections reveal a biphasic microstructure of α

(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.

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the number of cycles between the “measurement cycles”is not

constant, and because small differences in the minimum load

are not avoidable, ΔD

min

was normalized by the number of cycles

over which the change occurred and by the minimum load,

which gives us ΔD

min-norm

(Figure 9). An overall decrease in

ΔD

min-norm

over Nis seen for all indents (Figure 9a–c), however,

with high ﬂuctuations from cycle to cycle at the beginning of

loading, up to N=10. Some indents even present negative

ΔD

min-norm

values. Up to the 100th cycle, ΔD

min-norm

approaches

values between 10

4

and 3 10

4

nm μN

1

. Over the further

course of loading (10

≤N≤10

), an overall decrease in

ΔD

min-norm

with slight alterations below 2 10

6

nm μN

1

observed. Note, however, that the ΔD

min-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 (Figure 9d) highlight the pronounced differences

in the extent of the ﬂuctuations seen over the ﬁrst ten cycles.

Further, for each of the indents, ΔD

min-norm

approaches clearly

distinguishable levels over the further course of loading

(Figure 9e,f).

3.3. Cyclic Nanoindent Morphology

Nanoindent and pile-up size and morphology were evaluated

quantitatively based on the SPM images. Maximum indentation

depth after fatigue loading (D

max

at the maximum number of

cycles, N=10

), projected indent area (A

), pile-up area (A

p-u

pile-up volume (V

p-u

), and maximum pile-up height (h

p-u,max

)

are shown for all nanoindents in Table 2. The linear correlation

Figure 5. 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.

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Figure 6. Schematic representation of the crystallographic orientation of α

based on the EBSD–IPF maps: differences in the appearance of α

come from

their different orientations and therefore different sectioning to the surface.

Figure 7. Cyclic deformation curves, D

a,p

over N, for a) all indents and

b) typical examples, representing the different progression types (nanoind-

ents 1, 3, and 16).

Figure 8. a) D

min

max

over Nfor all indents. b) Examples of indents that

differ in their extent of curve progression.

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coefﬁcients (R) for the different parameters are summarized in

Table 3. The strongest correlations are seen between D

max

(N=10

) and A

(R=0.78), and between A

and A

p-u

(R=0.71), whereas the weakest correlations were found for

p-u

and h

p-u,max

(R=0.11) and for h

p-u,max

and V

p-u

(R=0.29).

Figure 10 displays SPM topography images and 3D surface

plots of typical nanoindents with different cyclic deformation

behavior, as described in Section 3.2. Indent 1 (Figure 10a) is

located in a surface region comprising α

-precipitates oriented

parallel to each other. It exhibits the largest indent and pile-up

sizes, regarding area (A

p-u

) and maximum height

p-u,max

), but an intermediate pile-up volume (V

p-u

). In the surface

region surrounding indent 3, the α

-precipitates are transversely

located (Figure 10b). This indent is about 21% smaller than indent

1(A

). Its lower deformation compared to indent 1 is also repre-

sented by smaller values of A

p-u

and h

p-u,max

.However,indent3

has the highest V

p-u

. Indent 16 (Figure 10c) is even smaller (A

)

with very little pile-up, as measured by V

p-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 speci-

men surface at positions with differences in the local distribution

and surface appearance of the α-phase (see HR-SEM insets in

Figure 11). 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 α

-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

-precipitates, oriented at different angles to each other. Both

indents C and D hit α

-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

Figure 9. Development of ΔD

min-norm

over a,d) 1 ≤N≤10

, b,e) 10

≤N≤10

and c,f) 10

≤N≤10

.(a–c) Curves for all nanoindents, and

(d–f) typical examples (nanoindents 1, 3, and 16) of the different behaviors observed.

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a smaller pile-up area. Indent D exhibits the smallest indent area

among the four indents, without signiﬁcant pile-up.

3.4. TEM Investigation

TEM (Figure 11) was used to evaluate the microstructure and the

dislocation density in volumes below the indents “A”to “D”.

Figure 11a shows nearly pure β-phase in the volume beneath

indent A, with only one α

-precipitate visible in the plane of

the TEM foil. A high dislocation density is observed along the

grain boundaries. Further, twins are seen near the indent

impression. For indent B (Figure 11b), a region with a high con-

tent of α

-precipitates is observed, and a high dislocation density

is seen mainly along grain or phase boundaries. A high number

of α

precipitates with different orientations is also revealed in

the volume directly beneath indent C (Figure 11c). In this case,

a higher dislocation density appears to be trapped between

adjacent, transversely positioned α-precipitates. In the volume

beneath indent D (Figure 11d), three α

-precipitates are oriented

parallel to each other, at an acute angle of about 135° to the

Table 2. Results of the quantitative analysis of indent and pile-up size:

projected indent area (A

), pile-up area (A

p-u

), pile-up volume (V

p-u

and maximum pile-up height (h

p-u,max

), together with the maximum

indentation depth after fatigue loading (D

max

(N=10

)) (n=indent

number). The smallest and greatest values for each parameter are

underlined and printed in bold, respectively.

Indent Pile-up

max

(N=10

p-u

p-u,max

p-u

[nm] [10

3

μm

] [10

3

μm

][μm] [10

4

μm

]

1272.72 4.77 11.25 0.19 1.91

2 231.28 3.81 6.81 0.15 1.82

3 223.43 3.76 7.11 0.01 2.84

4 222.48 4.18 6.79 0.11 1.47

5 217.56 3.39 6.51 0.12 0.36

6 213.44 3.61 7.47 0.14 0.88

7 214.1 3.84 10.77 0.10 1.48

8 205.19 3.57 10.62 0.11 1.25

9 204.76 2.96 2.52 0.13 0.36

10 196.16 2.91 5.27 0.07 0.77

11 208.03 3.14 9.85 0.08 0.83

12 197.49 2.67 4.07 0.15 0.60

13 209.54 3.91 9.68 0.06 1.12

14 213.55 3.73 7.41 0.1 0.86

15 211.97 3.12 5.28 0.12 0.41

16 198.81 3.01 4.12 0.09 0.54

17 195.63 3.19 6.48 0.09 0.31

18 200.66 3.38 6.14 0.13 0.92

19 194.13 2.49 3.10 0.06 0.30

20 220.36 3.74 6.53 0.08 0.70

21 212.55 4.19 7.14 0.13 0.96

22 217.21 3.14 3.98 0.11 0.63

23 216.72 3.82 6.84 0.12 0.89

24 210.74 3.08 5.57 0.13 0.86

Table 3. Linear correlation coefﬁcients (R) between the indent and pile-up

parameters given in Table 2. The two weakest values of Rare underlined,

and the two strongest values are given in bold font.

Parameter 1 Parameter 2 R

max

0.78

p-u

0.71

p-u,max

0.37

p-u

0.65

p-u

max

0.49

p-u,max

0.11

p-u

0.54

p-u,max

max

0.60

p-u

0.29

p-u

max

0.61

Figure 10. 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.

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surface. Around these particles, a high dislocation density is

observed.

4. Discussion

We used cyclic nanoindentation to characterize the inﬂuence of

the local microstructure on the cyclic deformation behavior of

Ti-5553 made by LPBF-M. Our results show a strong correlation

between the cyclic deformation data, the microstructure in the

volume beneath the nanoindents, and the fatigue-induced sur-

face morphology and dislocation structure. The orientation

and distribution of α

play a critical role in the cyclic deformation

mechanism of this metastable β-Ti alloy, as discussed in detail in

the following sections.

4.1. Microstructure

Through heat treatment of the as-printed specimens below the

β-transus temperature, we achieved an (αþβ)-microstructure,

with small α

-particles evenly distributed in the β-matrix, as to

be expected based on results reported for classically manufac-

tured Ti-5553.

[36,37]

The α

-phase fraction in our material

is in the lower range of values reached for the nonadditively-

manufactured materials (12% vs. 10%

[38]

to 26%

[39]

). Further,

in contrast to α

-chevrons,

[40]

we observe singular, acicular

Figure 11. a–d) Bright-ﬁeld TEM micrographs of the volumes beneath the indents “A”to “D”together with HRSEM 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.

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particles, which may be due to the fundamentally different

processing conditions before the heat treatment.

Over signiﬁcant regions of each β-grain, the α

acicular pre-

cipitates 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, our EBSD

measurements do not indicate an overall preferred crystal orien-

tation and they do not hint at an orientation relationship between

the α

-phase and the surrounding β-matrix. The latter result dif-

fers from the observation of others who reported that the α

phase has a Burgers orientation relationship with the β-matrix

in (αþβ)-solution annealed Ti-5553.

[41]

A likely explanation is

differences in the processing of conventional and additively man-

ufactured Ti-5553. The conventional process involves thermome-

chanical treatments before the actual heat treatment. Thus, a

lamellar (αþβ) microstructure is the starting point, while in

our case it is pure β, which has neither been plastically deformed

nor recrystallized.

4.2. Cyclic Deformation and Creep Behavior

We observe typical changes in hysteresis parameters (D

min

max

, and ΔD

min-norm

) over the number of cycles. The over-

all trend of the plastic displacement amplitude (Figure 7) over the

whole test reveals three subsequent stages of hardening, satura-

tion, and softening. Reports on the behavior of conventionally

processed Ti-5553 on the macroscale differ from our observation,

and they are not consistent: only cyclic softening, cyclic softening

followed by saturation, or cyclic hardening followed by softening

were observed.

[12,42,43]

Microstructural and compositional differ-

ences are one likely reason for the differences. All the reports

refer to conventionally processed metastable β-alloys, with differ-

ent compositions and heat treatments, compared to each other

and our alloy. Moreover, classical macrofatigue tests reveal an

average response over the microstructural constituents, whereas

we probe the local interactions of phases with the deformation

mechanisms (dislocation formation and movement, twinning),

without averaging. Thus, we extract the inﬂuence of local struc-

tural inhomogeneities on the cyclic deformation behavior, which

also explains the scatter between different indents (=regions)

and the ﬂuctuations we see in the progression of some hysteresis

parameters (see below).

The second parameter that we evaluated, ΔD

min-norm

decreases continuously and signiﬁcantly with increasing num-

bers of cycles, while D

min

max

increases steadily. Like D

ΔD

min-norm

exhibits signiﬁcant ﬂuctuations at the beginning of

the tests. Such ﬂuctuations are not seen for D

min

max

. 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 D

max

), D

min

max

, further hints at an overall

more elastic unloading behavior with ongoing cyclic deforma-

tion, which correlates with the saturation observed in the

progression of D

a,p

Our TEM investigations suggest that the interaction of dislo-

cations with α

-precipitates is the most important cyclic deforma-

tion mechanism inﬂuencing cyclic hardening and softening and

cyclic creep. The existence and orientation of the α

-precipitates

below and around the indents inﬂuence the formation of dislo-

cation structures. Most nanoindents exhibit alternating harden-

ing and softening with an overall trend for smaller plastic

displacement amplitudes (cyclic hardening) over the ﬁrst ten

cycles. Based on observations from macrofatigue tests,

[12,14]

we hypothesize that the repeated indentation activates multiple

slip systems, as well as interactions of dislocations with each

other and with nearby α

-precipitates, leading to hardening.

As for macrospecimens, 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 (αþβ)-Ti-5553

and (αþβ)-Ti-1023.

[12,14,42]

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≤10

, we observe only small further changes

in the plastic deformation amplitude, with some indents show-

ing overall hardening, and others overall softening. In this

“nearly saturation”state, therefore, either the described harden-

ing or the softening mechanisms are dominant. Further, with

ongoing loading, more dislocations can be activated and interact

with differently oriented α

-precipitates in the indented volume,

and the interaction volume expands, to a smaller or greater

extent, depending on the existence and orientation of α

-particles

nearby. Transversely placed α

-precipitates block the motion of

the dislocations more effectively than α

whose long axis is ori-

ented orthogonal to the surface, that is, parallel to the indentation

direction (compare Figure 11c,d). Thus, the interaction volume

can expand more in the latter case, overcoming possible disloca-

tion annihilation and strengthening the β-matrix. α

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 micro-

mechanisms.

[12]

Here, α

-precipitates in one sample deformed

to different strain levels depending on their orientation to the

loading direction.

For N≥10

, 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 α

-precipitates, together with the impingement of

α/βphase boundaries to dislocation movement.

[12]

Softening

is likely due to new dislocation arrangements occurring in α

and in the volume below the indents, thereby facilitating plastic

deformation.

[12,14]

Additionally, cyclic softening can be intensi-

ﬁed with dislocations relocating from regions of high dislocation

density to regions of low density.

[44]

Another mechanism is twin formation, shown by β-twins ini-

tiated at grain boundaries during cyclic nanoindentation (see red

arrow in Figure 11a). With increasing deformation, twin bound-

aries can progressively form and effectively block dislocation

movement.

Especially during the ﬁrst ten cycles, we observe relatively

large ﬂuctuations in the progressions of D

and ΔD

min-norm

over

N. 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 ΔD

min-norm

occasion-

ally even acquires negative values. These indicate that the

indenter is pushed up, instead of being pushed to the same

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or a bigger depth than in the cycle before. A similar behavior was

observed during cyclic nanoindentation of an Al–Si alloy.

[27]

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 heat-

ing and cooling in nearby regions may lead to quickly changing

states of thermal expansion and contraction.

[45,46]

An additional

cause can be modiﬁcations of the local microstructure due to the

β!αphase transformation during the heat treatment.

[47]

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 α

-precipitates (Figure S2,

Supporting Information). 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.

4.3. 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 (A

)

correlate with greater indent depths (D

max

), suggesting an overall

lower resistance to cyclic plastic deformation in the indented

volume. This is reﬂected by higher values of D

min

max

and

ΔD

min-norm

curves and pronounced cyclic softening for N≥10

(see, e.g., nanoindent 1). Correspondingly, the smallest A

and

max

values correlate with cyclic hardening and the lowest values

for ΔD

min-norm

(see, e.g., nanoindent 16).

The extent of pile-up is strongly inﬂuenced by the microstruc-

ture surrounding and below the indents, determining to what

extent plastic deformation is hindered in the volume below

the indent. For our 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., Figure 11c). Consequently, they cause a

decrease in the amount of plastic deformation.

[48]

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.

[27]

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 in β-grains (Figure 11a) and large pile-up happens above

the twin boundaries.

Cyclic indentations performed in regions with a relatively high

content of widely spaced α

-precipitates have intermediate sizes

and large pile-up volumes (e.g., nanoindent B, Figure 11b). 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 α

during the heat treatment deforms

the β-matrix, causing considerable stress and thus generating dis-

locations. During cyclic indentation, the high dislocation density

is released from the α

-precipitates that were directly encoun-

tered 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 leads to an intermediate cyclic plastic deforma-

tion state, as exempliﬁed by indent 3 in Figure 9f.

In comparison, small nanoindents usually arise from cyclic

indentation performed directly on α

-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 α

enables dislocation movement with-

out signiﬁcant restrictions, as exempliﬁed in Figure 11d. A large

Figure 12. Schematic representation of dislocation distribution and structures in the volume beneath cyclic nanoindents: a) dislocation allocation

affected by a low number of α

-precipitates oriented in a way such that a direct path for the dislocations toward deeper regions is available;

b) high content of differently oriented α

, 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.

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volume of material is plastically deformed and the dislocations

move deeper below the indent, if only a low amount or no α

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.

4.4. Summarizing Model Mechanism

Summarizing our results, we propose the interaction

mechanism shown schematically in Figure 12. Precipitates with

a preferred orientation where their long axes are parallel to the

indentation direction are less effective hindering dislocation slip

(Figure 12a) than precipitates that are not coaligned (Figure 12b).

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 α

results in more space for the dislocation

mobility and thus higher plastic deformation values and more

cyclic creep (Figure 12c).

5. Conclusion

We used cyclic nanoindentation to investigate local fatigue

processes in an LPBF-M β-metastable Ti-5553 alloy with a binary

(αþβ) microstructure. The α

-precipitates play an important role

in the local fatigue behavior. 1) Dislocation-based deformation

mechanisms are the main origin of cyclic softening, cyclic hard-

ening, and creep processes during cyclic nanoindentation. 2) A

strong correlation between the microstructure in the volume

beneath the nanoindents and the dislocation reaction was

observed: α

-phase orientation and distribution within the

β-grains signiﬁcantly contribute to the effectiveness of the precip-

itates as barriers to dislocation motion. High density, gathering,

and trapping of dislocations were observed at α/βinterfaces.

3) Pile-up occurrence and size are determined by the local plastic

deformability, which in turn is signiﬁcantly inﬂuenced by the

presence and orientation of α

-precipitates. 4) Our ﬁndings

differ from the results reported for the macrofatigue behavior

of (αþβ)-Ti-5553 alloy, as our testing approach considers the

inﬂuence of local structural inhomogeneities on the cyclic

deformation.

Concluding, our results highlight the high potential of cyclic

nanoindentation to elucidate the inﬂuence of the local micro-

structure on the cyclic deformation mechanisms of two-phase

alloys, such as the novel implant alloy investigated. The gained

understanding of the local interactions is the basis for improving

the fatigue performance of this alloy on the macrolevel.

Supporting Information

Supporting Information is available from the Wiley Online Library or from

the author.

Acknowledgements

The authors thank G. Gerlitzky and E. Uhlmann, Chair of Machine

Tools and Manufacturing Technology, TU Berlin, for providing the

Ti-5553 specimen. The authors further acknowledge M. Schaube for metal-

lographic preparation; R. Engelmayer and R. Meinke for technical support;

and A. Märten and M. Schmahl for support with the nanoindentation tech-

nique (all Chair of Materials Science & Engineering, TU Berlin). The

authors further thank I. Nigro (Dept. of Mechanical Engineering,

Politecnico di Milano, Italy) for her support with data collection and anal-

ysis. For providing the Python program to evaluate the nanoindentation

data, the authors thank K. Winkler, XPLORAYTION GmbH, Berlin, and

D. Hübler and R. Meinke, Chair of Materials Science & Engineering,

TU Berlin. The authors are further grateful to C. Fahrenson, C.

Günther, and S. Selve, Central Electron Microscopy Unit (ZELMI), TU

Berlin, for high-resolution scanning electron microscopy, focused ion

beam preparation, and transmission electron microscopy. The authors

acknowledge cofunding by the German Research Foundation, DFG, for

the nanoindenter (grant no. GZ INST 131/718-1).

Open Access funding enabled and organized by Projekt DEAL.

Conﬂict of Interest

The authors declare no conﬂict of interest.

Data Availability Statement

The data that support the ﬁndings of this study are available from the

corresponding author upon reasonable request.

Keywords

βTi-alloys, cyclic creep behavior, cyclic deformation behavior, cyclic

nanoindentation, implant materials

Received: July 18, 2023

Revised: September 2, 2023

Published online: September 20, 2023

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