Advanced Nanoindentation Testing for Studying Strain-Rate Sensitivity and Activation Volume
Advanced Nanoindentation Testing for Studying Strain-Rate Sensitivity and Activation Volume
VERENA MAIER-KIENER 0
KARSTEN DURST 0
0 1.-Department Physical Metallurgy and Materials Testing, Chair Physical Metallurgy and Metallic Materials, Montanuniversita ̈t Leoben , Roseggerstr. 12, 8700 Leoben , Austria. 2.-Physical Metallurgy, Technical University Darmstadt , Alarich-Weiss-Str. 2, 64287 Darmstadt, Germany. 3.-
Nanoindentation became a versatile tool for testing local mechanical properties beyond hardness and modulus. By adapting standard nanoindentation test methods, simple protocols capable of probing thermally activated deformation processes can be accomplished. Abrupt strain-rate changes within one indentation allow determining the strain-rate dependency of hardness at various indentation depths. For probing lower strain-rates and excluding thermal drift influences, long-term creep experiments can be performed by using the dynamic contact stiffness for determining the true contact area. From both procedures hardness and strain-rate, and consequently strain-rate sensitivity and activation volume can be reliably deducted within one indentation, permitting information on the locally acting thermally activated deformation mechanism. This review will first discuss various testing protocols including possible challenges and improvements. Second, it will focus on different examples showing the direct influence of crystal structure and/or microstructure on the underlying deformation behavior in pure and highly alloyed material systems.
The strain-rate sensitivity exponent and the
activation volume provide a fingerprint of the rate
controlling mechanisms during thermally activated
deformation. Over many years, nanoindentation
has been used for assessing local mechanical
properties such as hardness and modulus. To shed light
into temperature and rate dependent deformation
behavior, indentation creep tests were suggested,1
and Mayo and Nix2 reported superplastic
deformation behavior of Pb and Sn during indentation.
Especially for ultrafine-grained (UFG) and
nanocrystalline (NC) materials, nanoindentation
became a frequently used technique to determine
strain-rate sensitivity (SRS). Schwaiger et al.
published in 2003
a systematic comparison of
nanoindentation and tensile testing on NC-Ni, where both
techniques found a strongly grain size dependent
SRS. Within the last decade, many studies on
different materials classes, microstructures and
environmental influences were reported in
literature. Due to the limited amount of required
material, especially the severe plastic deformation
(SPD)4 and thin film5,6 communities used
nanoindentation to gain information on the microstructure
dependent deformation behavior of many
face-centered cubic (FCC),7–16 body-centered cubic
(BCC),17–21 and some hexagonal closed packed
(HCP)22–25 metals. Also the SRS of nanocomposites
such as Cu-Nb,26 nanoporous Cu and Au,27,28 or
nowadays high-entropy alloys (HEA)29,30 were
successfully investigated by nanoindentation. Besides
the successful determination of positive SRS, few
studies report experiments and results of negative
SRS.31,32 Moreover, many other investigations
dealing with various materials systems, such as
glasses32–34 or intermetallic phases,35 can also be
The general analyses of SRS probed by
nanoindentation is based on the conventional methods for
macroscopic materials testing.36,37 For more
detailed information, the reader is referred to
classic literature such as the textbook by Caillard.38
Generally, the parameter quantifying the SRS of a
material is m. It can be directly calculated from a
simple power-law relationship between plastic
stress r and applied strain-rate e_ (r ¼ e_m). For
nanoindentation experiments, correspondingly the
indentation stress, thus the hardness H, is used:
v ¼ kT
v ¼ C
Many macroscopic studies regarding thermally
activated deformation processes can be found in the
literature.43–45 Moreover, especially the comparison
between local and global behavior was intensively
used to get a better understanding of the acting
deformation mechanism in UFG-materials.46,47
Nevertheless, one major aspect concerning all these
studies are some experimental issues occurring
during long-term nanoindentation, mainly caused
by insuperably and omnipresent thermal drift.
In the following, two independent advanced
nanoindentation methodologies for a reliable
characterization of a strain-rate sensitive deformation
behavior will be discussed on the example of a
highly deformed austenitic steel (A220).
Furthermore, several examples will demonstrate the
significance of microstructure, crystal structure, and
solute content on the time dependent deformation
behavior and the underlying thermally activated
All shown indentation data and protocols were
carried out using a standard platform Nanoindenter
G200 (Keysight Tec, San Jose, USA). The
instrument was equipped with a Berkovich diamond tip
and a dynamic indentation unit (continuous
stiffness measurement—CSM), where a harmonic
oscillation is applied and the phase angles, the load, and
the displacement amplitude together with the
Sneddons equation48 are used to analyze the dynamic
contact stiffness. This allows the determination of
modulus and hardness continuously over
indentation depth. For further information on that
techniques and possible issues, the reader is referred to
the literature.49–51 At this point, it should be
mentioned that all presented approaches are not
limited to this manufacturer but can be readily
implemented on all available indentation systems
equipped with a dynamic indentation unit.
EXPERIMENTAL POSSIBILITIES AND
CHALLENGES: DEVELOPMENT OF
ADVANCED TESTING TECHNIQUES
The basic principle for probing thermally
activated deformation is to analyze the time dependent
material response during a stress- or strain-rate
controlled experiment, i.e., during strain-rate
changes, stress relaxation, or creep loading. In
nanoindentation, several testing possibilities are
available, such as constant strain-rate,52
indentation creep,53 or indentation relaxation54–56
experiments. In all cases, the indentation strain-rate is
the key parameter for a successful analysis. In the
following discussion, several examples with possible
issues are addressed, focusing mainly on pyramidal
indenter tips (constant indentation strain,
dependent on opening angle39) and only touching
Overview of Various Testing Options
The simplest but overall most frequently used
indentation method is the constant load rate (cLR)
indentation, where the indentation is conducted
with a constant loading rate to a preset load, mostly
without using any dynamic indentation technique.
Therefore, loading times, load levels, and number of
loading cycles can be adjusted by the user. Since in
1999 when Lucas and Oliver52 successfully showed
that a constant indentation strain-rate can be
accomplished by a proportional loading protocol, it
has been obvious that the applied indentation
strain-rate during each indentation segment is not
constant. According to Lucas and Oliver, the
indentation strain-rate can be described as follows (with h
indentation depth, P indentation load, and H
h_ 1 P_ H_
e_ ¼ h ¼ 2 P H
Thus, for a constant strain-rate (cSR) experiment
with pyramidal indenter tip geometry, a
proportional loading protocol according to Eq. 4 is applied.
This is commonly used in combination with a
dynamic indentation unit, which analyzes, based
on the dynamic stiffness and Sneddons equation,
the mechanical data continuously as a function of
indentation depth. Generally, the cSR is often used
in the literature, but it is also known that
macroscopic and local data were prone to differ in
between,9 mainly caused by thermal drift effects.
As shown in Refs. 32 and 57, cSR indentations with
a strain-rate of 0.001 s 1 might easily end up with
individual indentation times of more than 2 h, while
a cSR of 0.05 s 1 leads to a total time of 200 s.
Taking into account low thermal drift values of
0.05 nm/s, this might easily sum up in displacement
drift of more than 350 nm for low cSRs.
Applying Eq. 4 to cLR indentations shows32 that
during conventional protocols with several partial
loading–unloading cycles, constant loading times,
and variable load levels, the strain-rate at the last
point of each loading segment generally varies. A
constant strain-rate is only achieved if the loading
time to the different load levels remains constant
during the partial loading segments, moreover, fully
unloading the indenter tip after each loading
segment. The corresponding contact stiffness and thus
the mechanical properties will be evaluated after a
holding segment at a constant peak load from the
elastic unloading curve.58 This hold segment,
however, still leads to some minor changes between the
cSR and cLR methods due to strain-rate sensitive
deformation behavior as recently pointed out by
Leitner et al.51
Nanoindentation Strain-Rate Jump Tests
To overcome these issues, Maier et al.57 and
Alkorta et al.19 came up with an advanced
indentation protocol, where abrupt strain-rate changes
are applied during one single indentation
segment.57 These kind of test are state-of-the-art in
many macroscopic compression or tension tests, and
they were successfully adjusted to the small-scale
testing world. Within nanoindentation strain-rate
jump tests, the applied strain-rate/indentation
depth protocol can be adjusted individually,
allowing variable indentation depths for the abrupt
changes and different indentation strain-rates.
Every parameter can be chosen to the requirement
of the application and material.32 Additionally,
several drawbacks of the cSR-method can be
overcome with this one indentation method. For
example, the total indentation time is significantly
reduced because the initial indentation depth is
performed at large strain-rates and segments with
lower indentation strain-rates are performed at
larger depths, effectively reducing testing times.
Moreover, strain-rate jump tests also allow probing
the mechanical properties at one single location,
thus, e.g., the determination of SRS of individual
phases or microstructural heterogeneities.57 For
stable microstructural conditions, the evaluated
strain-rate sensitivity is independent of the
individual applied sequences of strain-rates and
indentations depths. Nevertheless, in the case of
unstable microstructures or during varying plastic
strain conditions, the experimental sequence might
slightly influence the evaluated materials
properties. Therefore, the SRS should be always discussed
as a function of hardness, since the rate sensitivity,
the microstrcutural length scales and the hardness
are closely related to each other.
In Fig. 1, the applied indentation protocol of a
nanoindentation strain-rate jump test (Fig. 1a) in
an austenitic steel A220 is shown together with the
corresponding load displacement curves (Fig. 1b) for
a coarse-grained (CG) and a highly deformed NC
microstructure. For both CG- and NC-conditions, a
significant dependency of the load–displacement
data on the applied strain-rate is found, with a
larger effect on the NC-condition. After each
SRjump, the curve goes through a transient region
adjusting correspondingly to the new applied SR.
Figure 1c shows the calculated Young’s modulus
and hardness data. The Young’s modulus is
independent of the applied SR, only showing more data
points in the regions with the lower indentation
strain-rates and thus longer testing times. After
each SR-jump, the hardness of the NC-condition
goes through a transient region until a new
SRdependent, but depth-independent hardness is
reached. Thereby, the hardness values of the first,
third, and fifth strain-rate segment (0.05 s 1)
exhibit about the same hardness values, indicating
that no indentation size effect or microstructure
changes are present.
To calculate a strain-rate sensitivity m or the
corresponding activation volume V , the hardness
difference according to the applied strain-rate
changes has to be determined, as shown in the
inset of Fig. 1c. These analysis have to be performed
in both a global manner (average values cSR
regions57) for materials showing a
depth-independent hardness and/or instantaneous (e.g., in the
presence of an ISE18) manner for materials with any
kind of indentation size effect as the size effect
would affect the global evaluation of m.
Nanoindentation Long-Term Creep Tests
Indentation creep or relaxation testing is another
common methodology to probe thermally activated
processes.54,59,60 In these tests, the indentation tip
is pushed into the material to a preset indentation
load or displacement, and afterwards, the load is
held constant for a preset creep time. To assess the
creep properties, the change in displacement has to
be recorded over a long time range.
The topic is highly discussed but at the same time
frequently used, and several analysis of
interpretation can be found in literature.59,61–67
Some main methodical aspects, especially while
comparing global and local data, have to be pointed
out and kept in mind:
During macroscopic creep testing, a constant
stress level is applied and a change in creep
strain and strain-rate are recorded. Pyramidal
indentation tips imply, however, a constant
representative strain to the material, which
varies in the plastic zone but is independent of
indentation depth. During constant load creep, a
change in the indentation hardness is found and
both applied stress and strain-rates are
changNevertheless, there the initial setting processes
in the plastic zone are hardly related to long time
To investigate the environmental influences during
nanoindentation creep testing at room temperature,
Maier et al.34 performed two individual constant
load indentations for 2 h at four different
indentation depths on the reference material fused quartz.
They found that the displacement signal for all
indentations showed some significant scatter that
obviously could not be explained with actual
mechanical properties. Further long-time
indentations of 24 h on FQ32 could prove that there is a
relation between indentation temperature evolution
and recorded displacement data.
These thermal drift related issues have already
been well known for a long time, although they are
still ignored in many works. In 1992, Weihs and
Pethica69 came up with a first dynamically
corrected long-term creep method. This method was
later further improved by Pethica and Asif70 and
Goldsby et al.71 All groups used a correction
technique based on the dynamic contact stiffness S
during both the loading and the constant load
segment. This contact stiffness is less prone to
thermal drift influences than the recorded
indentation displacement. While the displacement is
recorded far away from the indented area in the
gauge of the indenter head, the dynamically
recorded contact stiffness gives knowledge about
the true apparent contact at a frequency of several
Hertz during the experiment. Together with
Sneddon’s equation,48 this allows the determination of
the true contact area Ac:
S ¼ pp
Ac ¼ SCSM
Fig. 1. Nanoindentation strain-rate jump tests demonstrated on
CGand NC-A220 (austenitic steel), (a) preset indentation strain-rate
protocol with four independent abrupt strain-rate changes, (b)
corresponding load–displacement curves, and (c) hardness and
Young’s modulus continuously over indentation depth (inlay presents
a magnification of the third strain-rate jump for visualization of
instantaneous SRS determination).
ing. This makes a comparison with macroscopic
constant stress creep tests rather complicated.65
Furthermore, during uniaxial creep testing, a
steady-state secondary creep regime is analyzed,
which cannot be reached with pyramidal
The use of spherical indentation tips during
creep67 probes a varying indentations strain
with changing indentation depth due to the not
self-similarity of the sphere, which comes closer
to macroscopic creep regimes but still suffers
from the not uniform stress–strain fields
underneath the indenter.63
Long time creep experiments can suffer from
thermal drift issues; therefore, some groups only
apply creep experiments of less than a minute.68
Ac ¼ 4b2
As described in Ref. 10 in more detail, from the
know contact area, the drift corrected contact depth
hc can be evaluated. Together with the Oliver–
P , drift issues can be
Pharr equation h ¼ hc þ e SCSM
taken into account for both displacement h and
contact depth hc. Finally, according to Joslin and
Oliver,72 this corrected contact area Ac can be also
used to determine the hardness, simply based on
contact stiffness and a known modulus:
Hcorrected ¼ Ac ¼ p
4b2 E2R SP2
A long time creep indentation experiment is shown
in Fig. 2 on A220. Figure 2a shows the mechanical
data of the cSR loading segment, including load–
displacement behavior, hardness, and modulus. The
reduced modulus Er from the loading segment
serves at the same time as basis for the dynamic
correction,10 as shown in Eqs. 5 and 6.
After reaching the preset indentation depth, the
raw load is held constant at that level for a preset
time (Fig. 2b). The recorded displacement data as
well as the corrected data according to the method
suggested by Weihs and Pethica69 are plotted in the
lower part. Although the thermal drift was less than
0.05 nm/s, the recorded original displacement data
differs significantly to the corrected one. The
corrected hardness and displacement exhibit some
scatter, which is due to the use of the square of
the dynamic contact signal. Any noise in the
dynamic stiffness is thereby amplified. Using a
power law function to fit the data, this can be
circumvented and the resulting values can be
treated analytically. Further analysis concerns the
determination of equivalent creep/strain-rates
e_equivalent ¼ hh_ from the corrected indentation signal
or a transformation of the hardness into an
equivalent stress requivalent ¼ HcoCrrected.
Applying this dynamic correction protocol to
several data sets shows extensively that the
mentioned differences in displacement signal just
collapse to one reproducible curve, as demonstrated in
A further successful application for this method is
the investigation of SRS of thin films, where the
thickness of the investigated film is too low for other
kind of tests.5 Additionally, for high-temperature
testing, especially for systems with an apparent
thermal mismatch due to a nonheated indentation
tip, this method might act as a possibility to
overcome the drastic issues of thermal drift during
indentation.10 But also for indentation systems with
an independently heated tip and sample, a
reduction of thermal drift issues especially during the
indentation times was demonstrated for ufg-Au.34
Other Small-Scale Testing Techniques
and Expanding to Harsh Environmental
Since nanoindentation testing is a common tool to
probe mechanical properties and is not limited
anymore to ambient conditions, several methods
and protocols can also be used for high-temperature
testing. Results of both nanoindentation strain-rate
jump tests8,17,35,68 as well as nanoindentation creep
tests26–28,34 are reported over a wide range of
temperatures and strain-rates, shedding more light
on local thermally activated deformation processes.
High temperature testing, however, leads to further
challenges regarding thermal stability and testing
The effect of different strain regimes on thermally
activated deformation processes can be addressed
by spherical indentation. Spherical tips have no
self-similar geometry; thus, the implied strain
varies with indentation stress and contact radius.
Recently, Feldner et al.74 implemented strain-rate
jump tests with a spherical tip. Localized creep tests
with spherical tips are also reported in the
literature.75 Spherical indenter tips have thereby many
advantages but also one drawback, which is the
required spherical shape, which is only easily
achieved for diamond indenters. Nevertheless, at
this point, it is worth mentioning that especially the
calibration of the nonperfect, spherical tip shapes,
the controlling parameters of cSR indentation
experiments, as well as the correct analysis of the
data are still in an emerging state.
Besides nanoindentation, which always
represents a multiaxial loading situation, uniaxial
micropillar compression experiments can also be
used to examine thermally activated deformation
parameters on a very local scale, using, for example,
creep,76 jump,73 or stress relaxation tests.77 Wehrs
et al.12,78 demonstrated that the different protocols
for the same material lead to comparable results.
Finally, also bulge test techniques were modified
regarding a reliable testing of SRS.79
Quantification of Thermally Activated
With a reproducible long-term creep and
strainrate jump test, the underlying thermally activated
deformation mechanisms can be determined for a
wide variety of materials. For the analysis, the m
and V are determined according to Eqs. 1 and 2.
Correlating these material parameters, a unique
characteristic for each material, depending on
crystal structure and microstructure but also on
temperature and deformation history, can be evaluated.
Figure 3a shows a Norton plot for A220 and various
other materials, where the calculated displacement
creep rate is plotted over hardness during the
constant load segment. Figure 3b displays the
corresponding SRS exponent from creep tests and
strain-rate jump tests for several fcc and bcc metals
as well as a glass and a ceramic. Generally jump
and creep tests are in good accordance with each
other,32,68 and as described for macroscopic
compression tests,80 the SRS depends on the hardness
(which is for jump tests mostly neglected) since
during creep experiments, m is given by the slope of
the hardness versus equivalent strain-rate.
The mechanical properties shown in Fig. 3 spread
over two orders of magnitude in hardness, five orders
in strain-rate, and two orders in strain-rate
sensitivity. Consequently, a variety of material behavior and
underlying deformation mechanisms were tested. We
will at this point not go into all details, but only
summarize a few general findings. Reference
materials such as FQ or Sapphire exhibit little SRS but are
definitively not negligible, which should be kept in
mind when using them as athermal standards.
Mg17Al12 as a representative for an intermetallic
phase shows at RT almost no thermal activation.
Nevertheless, overcoming 150 C, it was reported that
the deformation behavior suddenly changes within a
narrow temperature interval to a highly
SR-dependent behavior.35 For pure Cr as a BCC metal,
underneath its critical temperature Tc, the SRS
increases with increasing grain size,17,18
approaching the single crystal condition due to the more
dominant thermally activated movement of screw
dislocations. The activation volume V , however, is
rather independent of the microstructure as the
kinkpair-mechanism prevails. In Al or Ni, which are
common examples for pure FCC-metals, the SRS
increases with grain refinement. Interestingly,
comparing pure Ni with a smaller grain size of 50 nm
compared to 300 nm for Al, the former exhibits a
lower rate sensitivity. Moreover, the hpt-A220 has a
grain size comparable to NC-Ni but lower values of
SRS. This indicates that simple thermally activated
dislocation annihilation at grain boundaries46 or
grain boundary sliding47 is not sufficient to explain
this behavior; rather, the interplay and detailed
atomistic details of the mechanisms involved should
be considered in the future. Initial studies on various
materials and microstructures indicate that besides
the dislocation-GB interaction, also solute-content
affected grain boundary motions might play a
dominant role, as focused in the next section.
Fig. 5. Influence of solid solution content and microstructure on SRS
of an equiatomic HEA with different microstructures (according to
Maier-Kiener et al.29 and reprinted with permission by Elsevier).
THERMALLY ACTIVATED PROCESSES
DEPENDING ON CRYSTAL STRUCTURE,
MICROSTRUCTURE, AND SOLUTE
Since strain-rate sensitive deformation behaviors
have manifold characteristics and also shed more
light into the individual underlying deformation
mechanism, in the following two subsections,
interesting materials systems will be discussed in closer
Zr-Based Materials—Influence of Crystal
Structure and Microstructure
Zr-based materials can be found in many different
structures and microstructures. Zr is common as a
base element in many BMGs, but also it is studied
as a pure metal in the nanocrystalline condition and
industrially used for example in cladding
During nanoindentation of BMGs, it is known
that the material shows strain-rate dependent
shear band formation.81,82 This can be also seen in
Fig. 4a for a Zr metallic glass (Zr-MG, for further
details, see Ref. 83), where the load–displacement
data show some serrated flow for low strain-rates.
Next to that, the data of CG- and a highly deformed
NC-condition are plotted. Corresponding Young’s
modulus and hardness data are shown in Fig. 4b
and c. While Zr-MG exhibits the highest hardness
but is independent of strain-rate, the modulus is, as
expected for BMGs, the lowest. CG-Zr has an hcp
crystal structure, which is also prone to show some
SRS in CG states due to the prismatic slip. This can
be clearly seen in the hardness data, leading to a
SRS of 0.023 and a corresponding V 13Æb3. The
transient behavior occurs smoothly, and the
materials require about 100-nm indentation depth until
a new steady deformation state is reached again.
SPD leads to a further change in the crystal
structure and to the formation of an x-phase,24
which can be also seen by the changed Young’s
modulus. The hardness significantly increases, but
the SRS stays the same. The characteristic of the
jump itself, however, has slightly changed to a more
abrupt distinct change in hardness with less
Highly Alloyed Materials—High-Entropy
HEA are a further intensively discussed
materials class. Besides their infinite number of
compositions and their unique mechanical properties, these
HEAs exhibit an interesting behavior regarding
their thermal activation.30,84,85 As already reported
for macroscopic tests,86 it was also shown during
nanoindentation of the CoCrMnNiFe-HEA that the
single crystalline like CH-condition exhibits a
surprisingly high SRS of around 0.01 for a single
crystalline FCC metal (Fig. 5c). A refinement of the
microstructure by HPT leads to a significant
increase in hardness and to an increasing SRS.
Corresponding behavior is seen in the activation
volume, where for all states, surprisingly low values
were calculated. According to these results, it can be
concluded that in these highly alloyed systems, a
further mechanism in the single crystalline state
might act since conventional pure fcc-metals are
expected to be strain-rate insensitive with
corresponding activation volumes around 1000Æb3.
Furthermore, with grain refinement, also the
jump characteristics change from a pronounced
transient behavior to a so-called yield point
phenomena, where the hardness first decreases
followed by a slight increase and a further decrease.
This is also found in other highly alloyed,
NCmaterials (such as NC-A220—Fig. 1c) and was also
reported previously by Van Petegem et al.87 for
NCNi-Fe alloys during transient tension testing.
Nevertheless, the overall underlying deformation
behavior is not fully understood yet and subject to
CONCLUSION AND OUTLOOK
Nanoindentation is a versatile tool for probing
thermally activated processes on a local scale. Simple
methodological adjustments lead to reliable testing
results paired with a minimization of testing time and
a reduction of number of indentations. This opens
further possibilities regarding a detailed investigation
of the properties of individual phases, heterogeneous
microstructures, and thin coatings. In this application
oriented review, different examples were presented to
demonstrate how SRS and activation volume can be
employed for a fundamental understanding of
governing deformation mechanisms but also how alloying
and solute content as well as microstructure influence
the thermally activated deformation behavior, in
terms of both the individual load–displacement
behavior as well as the overall hardness evaluation.
Generally it is stated that:
Distinct testing protocols have to be carefully
selected regarding reliability, application, and
Strain-rate jump tests allow the calculation of m
and V in a limited amount of time based on a
Dynamically corrected long-term creep
experiments can be used to probe low strain-rate
regions, but the achieved strain-rates depend
on the SRS of the material itself.
Crystal structure, microstructure, and testing
temperature significantly influence the
The absolute values of strain-rate sensitivity m
and activation volume V can be directly used to
identify the underlying rate-limiting deformation
mechanism. Moreover, the quantitative changes
upon, for example, different material
modifications, loading rates, or temperatures are
indicative of the interaction, competition, and change of
different contributing plasticity mechanisms.
Nonetheless, for an overall scale bridging and
comprehensive understanding of thermally
activated processes, macroscopic and microstructural
investigations are additionally recommended to
allow for a reliable investigation over several length
Open access funding provided by
Montanuniversita¨ t Leoben. The authors especially thank M.
Go¨ken and all colleagues at the Chair of General
Materials Properties (FAU Erlangen-Nu¨ rnberg) for
their support and collaboration during their time
there. V.M.-K. thanks A. Hohenwarter, A. Leitner,
B. Schuh (Montanuniversita¨ t Leoben), L. Kra¨ mer,
and O. Renk (O¨ AW-ESI) for providing some of the
here presented materials. R. Pippan, D. Kiener, and
H. Clemens are gratefully acknowledged for many
fruitful discussions and their support. Financial
support by the Austrian Federal Government
(837900) within the framework of the COMET
Funding Programme (MPPE, Project, A7.19) and by
the European Research Council under Grant
Number 340185 is appreciated (V.M.-K.).
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