Shap. Mem. Superelasticity
https://doi.org/10.1007/s40830-021-00356-9
TECHNICAL ARTICLE
Continuous Heating Dissolution and Continuous Cooling
Precipitation Diagrams of a Nickel-Titanium Shape Memory
Alloy
Christian Rowolt1 • Benjamin Milkereit1,2
Olaf Kessler1,2
•
Jette Broer1 • Armin Springer3
•
Received: 12 October 2021 / Revised: 9 November 2021 / Accepted: 10 November 2021
The Author(s) 2021
Abstract Binary NiTi alloys are the most common shape
memory alloys in medical applications, combining good
mechanical properties and high biocompatibility. In NiTi
alloys, the shape memory effect is caused by the transformation of an austenite phase to a martensite phase and
the reverse process. Transformation temperatures are
strongly influenced by the exact chemical composition of
the NiTi phase and the presence of precipitates in the
microstructure induced by thermo-mechanical treatment,
especially solution annealing and ageing. Isothermal time–
temperature precipitation diagrams can be found in the
literature. Cooling is frequently not considered, as water
quenching is typically assumed to be sufficient. To the best
of our knowledge, continuous heating dissolution (CHD)
and continuous cooling precipitation (CCP) diagrams do
not exist. Differential scanning calorimetry (DSC) is a
common method to analyse the austenite/martensite
transformation in shape memory alloys, but it has not yet
been used to analyse precipitation processes during continuous temperature changes. We have enabled DSC to
analyse dissolution and precipitation processes in situ
during heating as well as during cooling from the solution
annealing temperature. Results are presented as CHD and
& Benjamin Milkereit
benjamin.milkereit@uni-rostock.de
1
Chair of Materials Science, Faculty of Mechanical
Engineering and Marine Technology, University of Rostock,
Albert Einstein-Str. 2, 18059 Rostock, Germany
2
Competence Centre CALOR, Department Life, Light &
Matter, Faculty of Interdisciplinary Research, University of
Rostock, Albert-Einstein-Str. 25, 18059 Rostock, Germany
3
Electron Microscopic Centre, University Medical Centre
Rostock, Strempelstraße 14, 18057 Rostock, Germany
CCP diagrams, including information from microstructure
analysis and the associated changes in the austenite/martensite transformation temperatures.
Keywords NiTi Differential scanning calorimetry
(DSC) Continuous heating dissolution diagram (CHD)
Dissolution Continuous cooling precipitation diagram
(CCP) Precipitation
Introduction
NiTi alloys are the most common shape memory alloys in
medical applications, combining good mechanical properties and high biocompatibility [1–5]. In binary NiTi alloys,
the shape memory effect is characterised by the transformation of austenite (B2-structure) to martensite (B190 structure) and the reverse process [6]. The low-temperature
B190 martensitic phase forms during cooling of the hightemperature B2 austenite phase [7]. In the solution-annealed and quenched condition, the transformation takes
place in one step from B2 to B190 during subsequent
heating or cooling in the temperature range from - 100
to ? 100 C. The transformation temperatures are strongly
influenced by the exact chemical composition of the NiTi
phase and the presence of precipitates induced by thermomechanical treatments, especially solution annealing and
ageing. The formation of Ni4Ti3 precipitates during ageing
or dislocations due to thermo-mechanical treatment hinders
the formation of B190 . Depending on the size and density
as well as the nucleation site density, these precipitates
disturb the direct transformation from B2 to B190 . Beyond
that, a second martensitic phase called the intermediate Rphase exists [3]. It has been observed that in the solutionannealed, quenched, and aged stage, the transformation
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Shap. Mem. Superelasticity
takes place in two steps in the temperature range from 100 to ? 100 C. The initial transformation is from B2 to
R-phase, which transforms to B190 and the
reverse [3, 7, 8]. The precipitation sequence during
isothermal ageing treatment in NiTi alloys can be described
as follows:
a (supersaturated B2 NiTi) ! a1 + Ni4Ti3
! a2 + Ni3Ti2
! a3 + Ni3Ti
ð1Þ
a1, a2, and a3 are matrices with different Ni concentrations
[11].
Typically, ageing is done below 600 C in the twophase region NiTi and Ni3Ti according to the NiTi phase
diagram [10, 12]. During ageing at lower temperatures and
shorter ageing times, Ni4Ti3 appears in the microstructure,
while at higher temperatures and longer ageing times, the
equilibrium phase Ni3Ti will be precipitated. For intermediate temperatures and times, the metastable Ni3Ti2
forms. Besides, it is known that the precipitation process
and resulting transformation temperatures are also dependent on the heating rate to the solution annealing/ageing
temperature as well as the cooling rate from the solution
annealing/ageing temperature [8, 11, 13–16]. However,
currently, no continuous heating dissolution (CHD) and
continuous cooling precipitation (CCP) diagrams for NiTi
alloys exist, and only isothermal time–temperature transformation
diagrams
can
be
found
in
the
literature [1, 11, 17].
Differential scanning calorimetry (DSC) is a common
method to analyse the austenitic and martensitic transformation in shape memory alloys (see, for example,
Refs. [1, 7, 9, 10]), but it has not yet been used to analyse
the phase transformations according to the precipitation
sequence of Ni3Ti (Eq. 1). Recently, DSC has successfully
been used to establish CHD and CCP diagrams for several
alloying systems, like Al, Fe, Mg, or Ni [18–20]. In this
work and for the first time, the DSC method was applied to
the analysis of dissolution and precipitation kinetics during
continuous temperature changes in a binary NiTi alloy.
Materials and Methods
The investigated alloy was a Ni-rich binary NiTi shape
memory alloy. In the delivered condition, the material was
a Ø 6 mm cylindrical rod with an austenite finish temperature of - 5 C ± 10 K. The cylindrical rod was formed
by means of round forging from a cast block that was
manufactured using vacuum induction melting (VIM). The
present alloy was delivered with a nominal composition of
50.9 at.% Ni and 49.1 at.% Ti. Before the DSC experiments, all samples were solution annealed (1000 C for
123
30 min in the Setaram Labsys Evo DSC, 700 C for
30 min in the Perkin Elmer Pyris Diamond DSC/8500
DSC) and cooled at a rate of 0.1 K/s to create a defined
initial condition for subsequent experiments. The sample
dimensions and resulting sample masses for the three used
types of DSC devices are as follows:
Setaram Labsys Evo DSC: Ø 4.9 mm x 11.6 mm,
& 1420 mg.
Setaram
Sensys
Evo
DSC:
Ø 5 mm x 14.4 mm, & 1850 mg. Perkin Elmer Pyris
Diamond DSC/8500 DSC: Ø 6 mm x 1 mm, & 180 mg.
The applied scan rates were varied over a wide dynamic
range, from 0.01 to 5 K/s for the heating experiments and
from 0.01 to 1 K/s for the cooling experiments. Further
cooling experiments with 0.001 to 0.03 K/s rates were
performed for metallographic analysis. The DSC data
evaluation for the heating experiments was performed
according to Ref. [21] and for the cooling experiments
according to Ref. [22]. The experimental setup to perform
high-temperature DSC with solution annealing temperatures of 1000 C in a heat flux DSC Setaram Labsys EVO
is described in Ref. [20]. The technical ceramic Rescor960
with the same geometry as the sample was employed as an
appropriate reference material for the baseline measurement. DSC samples were packed in commercial ceramic
crucibles. Two sample measurements and one baseline
measurement were performed for each parameter setup.
For DSC experiments with temperatures up to 700 C,
the heat flux DSC Setaram Sensys Evo and the power
compensated Perkin Elmer Pyris Diamond DSC were used.
According to the Ni–Ti phase diagram, a temperature of
700 C is sufficiently high to anneal in the single phase
field of NiTi [12]. Later on, the DSC experiments will
prove this assumption. In this case, the stable a-titanium
alloy Ti5Al2.5Sn with the same geometry was used as a
reference material for maximum temperatures of 700 C.
To investigate the influence of quench-induced precipitation during cooling from the solution annealing temperature on the martensite transformation, the
determination of transformation temperatures from
austenite to martensite and the reverse was conducted.
Therefore, the standard test method described in
ASTM F2004-05 was adapted using a Perkin Elmer Pyris
Diamond
DSC
with
sample
dimensions
of
Ø 6 mm 9 1 mm. After solution annealing at 700 C for
30 min and cooling with different rates, a common ageing
treatment at 500 C for 20 min was performed in a protective nitrogen atmosphere. Afterwards, the temperature
range was set from - 80 to ? 100 C at controlled heating
and cooling rates of 0.05 K/s. Fig. 1 shows the applied
time/temperature profile for determination of the transformation temperatures. In the case of these experiments, one
sample was scanned multiple times in a series of varied
cooling rates after solution treatment. This experiment
Shap. Mem. Superelasticity
Fig. 1 Time/temperature profile for the determination of transformation temperatures after solution annealing at 700 C, varying the
cooling rate from the solution annealing temperature, and subsequent
ageing at 500 C for 20 min
The embedded and polished samples were mounted on the
SEM carrier with adhesive conductive carbon and aluminium tape (Co. PLANO, Wetzlar, Germany). The
acceleration voltage for the EDS analysis was set to 20 kV.
Figure 2 shows metallographic images obtained by SEM
(polished and unetched) in the initial condition of the NiTi.
In the cross section in Fig. 2a and b, many small black
particles can be observed in the grey matrix. In Fig. 2c and
d, images are displayed in the longitudinal direction where
aligned particles are evident. EDS proves that the inclusions show a high concentration of titanium and carbon.
Furthermore, the nickel content is significantly lower than
in the surrounding matrix. The large size (up to a few
hundred micrometres) and the high density of those particles correlate with non-metallic inclusions like titanium
carbides and oxides, which were also observed in
Ref. [23–25]. According to the literature, these inclusions
can be traced back to the primary shaping process.
Results and Discussion
series focussed on the determination of variations in the
characteristic transformation temperatures. No baseline
measurements were done in this special case. This is to
allow programming of the series in one experiment-program, and as the raw heat flow was found to be sufficiently
precise to allow evaluation of relative variations in the
transformation temperatures.
To analyse the changes in the microstructure related to
the reactions detected in DSC, optical microscopy (OM) as
well as scanning electron microscopy (SEM) and energy
dispersive X-ray spectroscopy (EDS) were performed. The
heat-treated samples were cold-embedded in epoxy resin
and then mechanically ground and polished with water-free
lubricants. The final polishing was done with a 0.05 lm
oxide suspension. The polished samples were rinsed and
etched as per following the etching agent: 62.5 ml water
H2O, 18.75 g potassium metabisulfite K2S2O5, 12.5 ml
hydrochloric acid HCl, and 1.5 g ammonium hydrogen
fluoride NH4HF2. Microstructural images were acquired by
optical microscopy using a LEICA DMI5000 M, Co. Leica Microsystems GmbH, Wetzlar, Germany.
The SEM samples were analysed by a field emission
SEM (MERLINVP Compact, Co. Zeiss, Oberkochen,
Germany) equipped with an EDS detector (XFlash 6/30,
Co. Bruker, Berlin, Germany) and analysis software
(Quantax400, Co. Bruker, Berlin, Germany). SEM-secondary electron (SEM-SE) images were obtained using a
high efficiency Everhart–Thornley-type HE-SE detector at
5 kV acceleration voltage. Representative areas of the
samples were analysed and mapped to determine the elemental distribution based on the EDS-spectra data by the
QUANTAX ESPRIT Microanalysis software (version 2.0).
Heating Experiments
DSC curves during continuous heating using a broad range
of heating rates are shown in Fig. 3. For a better understanding, the curves have been offset downwards from the
slowest rate on top to the fastest at the bottom. Horizontal
dotted lines correspond to the zero level for the curves.
Upward deviation of the DSC curve from the zero level
indicates domination of endothermal reactions, whilst
downward deviations of the DSC curve from the zero level
indicate domination of exothermal reactions. One must be
aware that the (partly opposing) reactions might superimpose and DSC is just measuring the sum of all heat effects
at present. Three different DSC devices were used, as
shown by the black, red, and blue curves. In general, there
is a good match between the three types of DSC devices
used to cover a wide range of heating rates as well as high
temperatures.
The DSC-heating curves of 0.01 K/s show an exothermal precipitation reaction between 230 and 430 C with a
distinct peak at 300 C (peak a). Additionally, an overlapping precipitation reaction appearing as a shoulder at
350 C can be seen (peak b). This holds for rates up to
0.1 K/s, while the shoulder is no longer seen at higher
heating rates. It is believed that in the context of these
reactions, parts of the precipitation sequence of Ni3Ti
(Eq. 1) occur. After crossing the zero level at 430 C, a
subsequent endothermal dissolution reaction with a peak at
580 C can be observed (peak C). It can be assumed that
previously precipitated particles will be dissolved during
this endothermal reaction. It can be seen that the
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Shap. Mem. Superelasticity
Fig. 2 Metallographic images
obtained by SEM in the initial
condition (polished and
unetched). a and b cross section,
c and d longitudinal direction
Fig. 3 Selected continuous DSC-heating curves using three different
devices: Setaram Labsys Evo (black curves), Setaram Sensys Evo
(red curves), and Perkin Elmer Pyris 8500 (blue curves) (Color
figure online)
dissolution reaction is finished at 600 C and solution
annealing above 600 C is sufficient to dissolve the
alloying elements. A further increase in the temperature
123
results in a horizontal progress of the DSC curve at zero
level, indicating a reaction free zone. With increasing
heating rate, the precipitation and dissolution reactions and
corresponding DSC peaks are significantly shifted to higher
temperatures (indicated by the vertical grey spline). In the
same order, they are increasingly suppressed. For the
fastest investigated heating rate of 5 K/s, the peak of the
exothermal precipitation reaction can be seen at 530 C,
which means a shift of about 230 K in comparison to
heating at 0.01 K/s. The peak of the endothermal dissolution reaction shifts from 580 C at 0.01 K/s to 620 C at
5 K/s. The suppression of ongoing reactions during heating
can firstly be seen by the disappearance of the shoulder
during the exothermal reaction for heating rates higher than
0.1 K/s as well as by the reduction of areas under the exoand endothermal peaks.
Figure 4 shows the resulting continuous heating dissolution (CHD) diagram for the specific batch and initial
condition of the NiTi alloy 50.9 at.% Ni and 49.1 at.% Ti
are investigated. Therefore, start and end temperatures of
characteristic precipitation and dissolution reactions were
evaluated and plotted in a temperature/time diagram. From
the CHD diagram, during heating to typical ageing temperatures, distinct precipitation and dissolution reactions
occur. This means that depending on the considered temperature and heating rate, certain stages in the precipitation
sequence of Ni3Ti (Eq. 1) can be reached prior to the actual
ageing treatment. This aspect must be considered when
selecting ageing temperatures and times as premature
Shap. Mem. Superelasticity
Fig. 4 Continuous heating dissolution (CHD) diagram of
50.9 at.% Ni and 49.1 at.% Ti alloy [initial condition: solution
annealed (1000 C for 30 min in Setaram Labsys Evo DSC, 700 C
for 30 min in Perkin Elmer Pyris Diamond DSC/8500 DSC) and
cooled at 0.1 K/s]
precipitation of nickel rich particles can significantly
influence on the transformation temperatures. In consequence different heating rates, especially slower heating
rates, can lead to different transformation temperatures
under otherwise identical ageing conditions. After running
through the different areas of precipitation and dissolution,
the heating rate specific solvus temperature is reached.
Above this temperature, all major alloying elements are in
a solid solution.
Cooling Experiments
DSC curves from continuous cooling using a broad range
of cooling rates are shown in Fig. 5. For a better understanding, the curves have been offset downwards from the
slowest rate on top to the fastest at the bottom. Horizontal
dotted lines correspond to the zero level for the curves.
Upward deviation of the DSC curve from the zero level
indicates exothermal reactions. During cooling, only
exothermic precipitation is expected; thus, the issue of
superimposed reactions is less critical for curve interpretation. Three different DSC devices were used, which are
shown by the black, red, and blue curves. Also, during
cooling, there is a good match between the three types of
DSC devices used to cover a wide range of cooling rates as
well as high-temperatures. DSC curves after solution
annealing at 700 and 1000 C are nearly identical. This
agrees with the CHD diagram in Fig. 4.
Considering a cooling rate of 0.01 K/s, one obvious
exothermal precipitation reaction occurs between 420 and
310 C. With increasing cooling rate, the precipitation
reaction is shifted to lower temperatures and gets
Fig. 5 Selected continuous DSC-cooling curves using three different
devices: Setaram Labsys Evo DSC (black curves), Setaram Sensys
Evo DSC (red curves), and Perkin Elmer Pyris Diamond DSC (blue
curves) (Color figure online)
increasingly suppressed. The peak temperature for 0.01 K/s
is 395 C, while it is shifted to 275 C at 1 K/s. However,
though the DSC curve of 0.01 K/s is quite noisy, another
weak, preceding precipitation peak might be present at
520 C. This assumption is supported by the shoulder seen
on the high-temperature side of the subsequent exothermic
peak at 0.03 K/s and 0.05 K/s. It can be assumed that this
reaction at higher temperatures would be even more pronounced at slower cooling as considered here.
Figure 6 shows the reduction of the total specific precipitation enthalpy with increasing cooling rate. The
specific precipitation enthalpy can be calculated by integrating the peak area. In this case, the highest specific
precipitation enthalpy is 3.6 J/g for the slowest cooling rate
of 0.01 K/s. This value is nearly constant for the next two
faster cooling rates. For 0.1 K/s, the specific precipitation
enthalpy drops down to 1.7 J/g and decreases further with
the increase in the cooling rate. For the highest investigated
cooling rate of 1 K/s, a precipitation enthalpy of just 0.3 J/
g remains, which is already close to the detection limit of
DSC [18, 22]. From this, it can be concluded that the
critical cooling rate to completely suppress quench-induced
precipitation is just slightly higher than 1 K/s (considering
a logarithmic cooling rate increment).
The start and end temperatures of the precipitation
reaction during cooling from 1000 C, respectively, to
700 C for 30 min were evaluated. The corresponding
123
Shap. Mem. Superelasticity
result in multiple overlapping precipitation peaks regarding
the precipitation sequence towards Ni3Ti. This behaviour is
well known from other alloying systems, for example, on
the basis of Al, Fe, or Mg [18, 19].
Transformation Behaviour after Cooling
with Different Rates
Fig. 6 Specific precipitation enthalpy and transformation temperatures after additional ageing at 500 C for 20 min as a function of
cooling rate (Color figure online)
time–temperature data points are plotted in Fig. 7 showing
the CCP diagram for the specific batch of the NiTi alloy
50.9 at.% Ni and 49.1 at.% Ti. The two fastest cooling
rates were performed starting from a temperature of
700 C; nevertheless, these two were integrated into the
same diagram. The critical cooling rate is estimated to be
3 K/s, as the specific precipitation enthalpy for 1 K/s is
almost zero (see Fig. 6). The dashed line proposes the
extrapolated start and end temperature for quench-induced
precipitation limited by the critical cooling rate at 3 K/s.
It can be assumed that quench-induced precipitation of
phases like Ni4Ti3, Ni3Ti2, or Ni3Ti from the precipitation
sequence towards Ni3Ti (Eq. 1) is observed here. As only
one clear exothermal precipitation peak occurs at a cooling
rate of 0.01 K/s, it is expected that slower cooling would
Fig. 7 Continuous cooling precipitation
50.9 at.% Ni and 49.1 at.% Ti alloy
123
(CCP) diagram of
The influence of the applied cooling rate variation
(0.01–5 K/s) from the solution annealing temperature on
the transformation behaviour between - 80 and ? 100 C
after additional ageing is shown in Fig. 8. The cooling step
including the endothermal transformation from the hightemperature austenite (B2) phase to the low-temperature
martensite (B190 ) phase is shown in Fig. 8a. During cooling of the sample previously cooled at 0.01 K/s from the
solution annealing temperature, at least three distinct peaks
can be found. During those peaks, a sequence of transformations occurs depending on the specific ratio of Ni/Ti and
the size, amount, and location of precipitates and dislocations [7–9]. It can be seen from the diagram that faster
cooling results in a slight shift of the transformation peak 1
to higher temperatures and the disappearance of the
transformation peak 2 for cooling rates higher than 0.3 K/s.
In addition, the transformation peak 3 significantly shifts to
lower temperatures with increasing cooling rate from 18 C at 0.01 K/s to - 38 C at 3 K/s. Comparing the
three fastest investigated cooling rates (1–5 K/s), no differences can be recognised for the transformation behaviour from austenite to martensite. This agrees with the
finding that the critical cooling rate is just slightly higher
than 1 K/s (Fig. 7).
In Fig. 8b, the subsequent heating from - 80 to ?
100 C is shown. During heating of the low-temperature
martensite (B190 ) phase, two distinct exothermal transformation peaks can be seen, resulting in the formation of the
high-temperature austenite (B2) phase. Peak 4 dominates
the transformation during heating. The transformation
peak 5 only shows up as a shoulder at transformation
peak 4. At the temperature at which the overall reaction
expires and adopts the course of the imaginary baseline, the
transformation is complete and the austenite finish temperature Af has been determined. Af has been evaluated
and summarised in Fig. 6. For a prior cooling rate of
0.01 K/s from the solution annealing temperature and
additional ageing, the Af temperature is determined to be
55.1 C. For the next faster cooling of 0.03 K/s, the Af
temperature drops to 48.6 C, staying nearly constant for
further increases in the cooling rate. Although the Af
temperature is nearly the same, the transformation behaviour differs. The shoulder from peak 5 gets increasingly
suppressed and shifted to lower temperatures.
Shap. Mem. Superelasticity
Fig. 8 Influence of the variation of cooling rate from the solution annealing temperature 700 C for 30 min on the transformation behaviour after
additional ageing at 500 C for 20 min during slow cooling (a) and heating (b) between -80 C and ?100 C (Color figure online)
In the same manner, we have evaluated the martensite
finish Mf and austenite start As temperatures (Fig. 8).
Increasing transformation temperatures Mf, As, and Af with
a lower cooling rate (Fig. 6) can be explained by the preceding precipitation from the solid solution during cooling.
Transformation temperatures rise with decreasing Ni content of the NiTi phase [6, 26], and the Ni content itself is
reduced by the precipitation sequence (Eq. 1). Thus,
stronger precipitation during cooling causes increasing
austenite/martensite transformation temperatures.
Microstructures after Cooling with Different Rates
The metallographic images from optical microscopy of the
investigated NiTi alloy are shown in Fig. 9. These samples
were continuously cooled from 700 C for 30 min to room
Fig. 9 Microstructure
development of NiTi after
varying the cooling rate from
700 C for 30 min in an etched
condition
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Shap. Mem. Superelasticity
temperature with varying cooling rates from 0.001 to
0.03 K/s to promote quench-induced precipitation. No
ageing treatment was added. In this state, transformation
from austenite to martensite starts significantly below room
temperature, i.e. the analysed microstructure consists
mainly of austenite. Etching reveals the austenite grains.
Beside the austenitic matrix, obviously a high density of
bright polygonal low aspect ratio particles of a few
micrometres in diameter exists randomly distributed on all
of the images. The size of a few micrometres and the
distribution of those particles correlate with non-metallic
inclusions like titanium carbides and oxides, which have
also been observed in Fig. 2 as well as by [23–25]. Hence,
these particles appear independent of the cooling rate.
After the slowest investigated cooling rate of 0.001 K/s,
at least two types of quench-induced precipitates are
observed. One type appears as elongated, bright-grey precipitates without a specific shape and orientation and a
length of about 20 lm. It seems that around these particles,
precipitate-free zones can be seen. That is, a large proportion of the alloying elements have diffused into these
large precipitates. At higher cooling rates, these precipitates are no longer visible, i.e. their formation during
cooling is suppressed. As a second type, many blue
coloured needle-shaped precipitates can be seen at
0.001 K/s, ranging in length from up to 20 lm down to a
few lm. With a thickness of only a few hundreds of nm,
they exhibit a high aspect ratio. These particles appear in
distinct orientations. Since the material under study is a
binary alloy of Ni and Ti, these precipitates are assumed to
come from different stages of the precipitation sequence
Ni3Ti (Eq. 1). However, within the scope of this work, we
need to keep the open question of which exact phases are
precipitated. Certainly, with increasing cooling rate, these
needles become smaller and less in number. After cooling
at 0.03 K/s, no quench-induced precipitates can be seen
using optical microscopy. As the DSC-signal clearly indicates an exothermal precipitation reaction during cooling at
this rate (Fig. 5), further particles are probably formed at
the nanoscale, which need further analysis by transmission
electron microscopy (TEM) in the future.
Conclusion
In this work, the successful application of a sophisticated,
systematic differential scanning calorimetry (DSC) in situ
analysis of dissolution and precipitation processes during
continuous heating and cooling of a Ni-rich binary NiTi
shape memory alloy has been shown for the first time.
From the systematic DSC analysis combined with optical
metallography, we conclude the following:
123
•
•
•
•
•
•
Quench-induced precipitation during cooling from the
solution annealing temperature was detected by DSC
and confirmed by microstructural analysis. The size and
number of precipitates decrease with increasing cooling
rate.
A critical cooling rate for quenching from the solution
annealing temperature of about 3 K/s has been
determined.
The impact of existing (quench-induced) precipitates
on the transformation temperatures has systematically
been analysed.
Pre-existing precipitates, i.e. nickel rich particles, can
significantly influence on the transformation temperatures. In consequence, different heating rates, especially
slower ones, can lead to different transformation
temperatures under otherwise identical ageing
conditions.
It can be further concluded that quench-induced
precipitation (QIP) also strongly influences the transformation behaviour of austenite to martensite and the
reverse. Specifically, with increasing suppression of
QIP, the martensite to austenite transformation temperature decreases. E.g. after cooling at various rates from
solution treatment and additional ageing (500 C
20 min), the martensite finish temperature during
cooling at 0.05 K/s from 100 C shifts from - 21 C
(preceding cooling at 0.01 K/s) to about - 41 C after
a preceding cooling at 5 K/s. A continuous heating
dissolution (CHD) and a continuous cooling precipitation (CCP) diagram for NiTi have been established.
These diagram can be used for the selection of heating
and cooling rates in processing NiTi components.
Future work needs to analyse the nature of quenchinduced precipitation on the micro and nanoscale.
Acknowledgements Financial support by the German Federal Ministry of Education and Research (BMBF) within the project
RESPONSE ‘‘Partnership for Innovation in Implant Technology’’
(Grant Number FKZ 03ZZ0928D) is gratefully acknowledged.
Author Contributions CR, BM, and OK designed the experiments.
CR performed and evaluated the experiments. CR wrote the manuscript. JB and AS performed the SEM analysis. All authors joined the
discussion and approved the final version of the paper.
Funding Open Access funding enabled and organized by Projekt
DEAL.
Declarations
Conflict of interest The authors declare that they have no conflict of
interest.
Open Access This article is licensed under a Creative Commons
Attribution 4.0 International License, which permits use, sharing,
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