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Ni-Ti alloys are considered to be very important shape memory alloys with a wide application area including, e.g., biomaterials, actuators, couplings, and components in automotive, aerospace, and robotics industries. In this study, the NiTi46 (wt.%) alloy was prepared by a combination of self-propagating high-temperature synthesis, milling, and spark plasma sintering consolidation at three various temperatures. The compacted samples were subsequently heat-treated at temperatures between 400 °C and 900 °C with the following quenching in water or slow cooling in a closed furnace. The influence of the consolidation temperature and regime of heat treatment on the microstructure, mechanical properties, and temperatures of phase transformation was evaluated. The results demonstrate the brittle behaviour of the samples directly after spark plasma sintering at all temperatures by the compressive test and no transformation temperatures at differential scanning calorimetry curves. The biggest improvement of mechanical properties, which was mainly a ductility enhancement, was achieved by heat treatment at 700 °C. Slow cooling has to be recommended in order to obtain the shape memory properties.
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materials
Article
Influence of Heat Treatment on Microstructure and
Properties of NiTi46 Alloy Consolidated by Spark
Plasma Sintering
Pavel Salvetr 1,* , Jaromír Dlouhý1, Andrea Školáková2, Filip Pr˚uša 2, Pavel Novák2,
Miroslav Karlík3,4 and Petr Haušild 3
1COMTES FHT, Prumyslova 995, 334 41 Dobrany, Czech Republic; jaromir.dlouhy@comtesfht.cz
2Department of Metals and Corrosion Engineering, University of Chemistry and Technology, Technická5,
166 28 Prague 6, Czech Republic; skolakoa@vscht.cz (A.Š.); prusaf@vscht.cz (F.P.); panovak@vscht.cz (P.N.)
3
Faculty of Nuclear Sciences and Physical Engineering, Department of Materials, Czech Technical University
in Prague, Trojanova 13, 120 00 Prague 2, Czech Republic; miroslav.karlik@fjfi.cvut.cz (M.K.);
petr.hausild@fjfi.cvut.cz (P.H.)
4
Faculty of Mathematics and Physics, Department of Physics of Materials, Charles University, Ke Karlovu 5,
121 16 Prague 2, Czech Republic
*Correspondence: pavel.salvetr@comtesfht.cz; Tel.: +420-220-44-4441
Received: 12 November 2019; Accepted: 3 December 2019; Published: 6 December 2019


Abstract:
Ni-Ti alloys are considered to be very important shape memory alloys with a wide
application area including, e.g., biomaterials, actuators, couplings, and components in automotive,
aerospace, and robotics industries. In this study, the NiTi46 (wt.%) alloy was prepared by a
combination of self-propagating high-temperature synthesis, milling, and spark plasma sintering
consolidation at three various temperatures. The compacted samples were subsequently heat-treated
at temperatures between 400
C and 900
C with the following quenching in water or slow cooling in
a closed furnace. The influence of the consolidation temperature and regime of heat treatment on
the microstructure, mechanical properties, and temperatures of phase transformation was evaluated.
The results demonstrate the brittle behaviour of the samples directly after spark plasma sintering at
all temperatures by the compressive test and no transformation temperatures at dierential scanning
calorimetry curves. The biggest improvement of mechanical properties, which was mainly a ductility
enhancement, was achieved by heat treatment at 700
C. Slow cooling has to be recommended in
order to obtain the shape memory properties.
Keywords:
Ni-Ti alloy; self-propagating high-temperature synthesis; spark plasma sintering; aging;
compressive test; hardness; shape memory
1. Introduction
The Ni-Ti alloys named NiTinol (derived from nickel, titanium, and laboratory of discovery-Naval
Ordnance Laboratory) are well-known shape memory alloys with good mechanical properties and
high corrosion resistance, which enables usage as implants, medical devices, and other applications
as biomaterials [
1
]. The shape memory eects occur due to the reversible solid-state transformation
between the high-temperature and low-temperature phases. The high-temperature phase is called
austenite with a high-symmetry structure-ordered body-centered cubic phase B2 (CsCl). The
low-temperature and low-symmetry phase is called martensite with monoclinic B19
0
lattice. Reversible
strains at about 8% of the initial length are enabled due to the reversible phase transformation as
well [
1
]. To achieve the desired shape memory eects-transformation temperatures, it is necessary to
be careful about a chemical composition. The transformation temperatures are very sensitive for the
Materials 2019,12, 4075; doi:10.3390/ma12244075 www.mdpi.com/journal/materials
Materials 2019,12, 4075 2 of 17
nickel-titanium ratio and an increase in the content of nickel by 0.1 at a percentage causes a change of
the transformation temperature A
F
(austenite finish) up to 10
C [
2
,
3
]. This fact makes the production
of the Ni-Ti alloys more dicult because titanium and, subsequently, NiTi alloys have high anity
to oxygen from atmosphere and carbon from the melting crucible during vacuum induction melting
(VIM) [4,5].
Spark plasma sintering (SPS) is a modern consolidation process, which is suitable for various
materials such as ceramics and metals including many intermetallic systems (e.g., Ni-Ti and Fe-Al
alloys [
6
9
]). The high heating rate and shortness of the whole process enable the use of SPS for
the consolidation of nanocrystalline materials as well [
10
12
]. The process is based on the sintering
of powder under the simultaneous influence of high electric current (direct or pulsed) and uniaxial
pressure. The Joule heat is generated by passing the current through the graphite punch and die
and between the powder’s particles. The high heating rate was described as the route to reduce the
amount of undesirable secondary phase such as the Ti
2
Ni by the self-propagating high-temperature
synthesis (SHS) [
13
]. Thus, the SPS process was examined as a heating source for initiating the SHS
reaction between nickel and titanium elemental powders in the previous paper. However, the SPS
process seems to be inapplicable for the initiation of the SHS reaction because the strongest increase of
the temperature occurs on the surface of the particles. The formed intermetallic layers act further as
diusion barriers and separate unreacted nickel from titanium [
14
]. Therefore, the pre-alloyed Ni-Ti
powder after mechanical alloying is usually sintered by SPS. This process can produce the highly
dense NiTi materials [
7
] or the porous structure depending on the addition of the space holder (e.g.,
NH4HCO3) [15,16].
Heat treatment of the Ni-Ti alloys has a crucial eect on the properties of the samples. The
parameters of heat treatment influence the microstructure, internal stresses, precipitation, shape
memory, and mechanical properties [
17
]. The Ni-Ti alloys undergo homogenizing treatment at the
temperature of about 800–1050
C up to several hours of duration, which is followed by water quenching
to get a homogeneous microstructure without the Ni-rich precipitates [
17
22
]. The second step of heat
treatment represents the aging treatment, which usually occurs between 300
C and 800
C [
23
,
24
]. The
grade of aging (density and size of the Ni
4
Ti
3
precipitates) depends on the temperature and time. The
precipitation process starts with the metastable Ni
4
Ti
3
phase, which is transformed into the metastable
Ni
3
Ti
2
phase and the stable Ni
3
Ti [
25
]. The precipitation process is accompanied by the hardness
changes during aging due to formation, coarsening, and decomposition of the Ni
4
Ti
3
and Ni
3
Ti
2
phases. The addition of Al enhances the microstructural and hardness stability of Ni-rich Ni-Ti alloys
until 500–600
C [
17
,
26
]. The results of phase formation, stability, and transformation Ni
4
Ti
3
Ni
3
Ti
2
Ni
3
Ti during aging are presented by a time-temperature-transformation (T-T-T) diagram [
17
,
25
,
27
].
In this work, the NiTi46 (wt.%) alloy was processed by a combination of SHS, milling in a vibratory
mill, and SPS consolidation at three temperatures to get fully dense materials. The prepared samples
underwent the heat treatment in temperatures ranging from 400 to 900
C. The characterization of
samples was focused on phase composition, observing the transformation temperatures, and changing
the mechanical properties depending on the heat treatment regime.
2. Materials and Methods
The metallic powders with the following particle sizes and purities were used as starting material
for the NiTi46 (wt.%) alloy: nickel (particle size <150
µ
m, 99.9 wt.% purity, Sigma-Aldrich, St. Louis,
MO, USA) and titanium (particle size <44
µ
m, 99.5 wt.% purity, STREM CHEMICALS, Newburyport,
MA, USA). The powders were mixed manually corresponding to the chemical composition of the
NiTi46 powder mixture, which was uniaxially compressed at room temperature to cylindrical green
bodies of 12 mm in diameter at a pressure of 450 MPa for 5 min using LabTest 5.250SP1-VM universal
loading machine (Labortech, Opava, Czech Republic). The SHS reaction of the pressed powder mixture
was carried out in the fused silica ampoules evacuated to 10
2
Pa and sealed, which were placed in the
preheated region to 1100
C electric resistance furnace. The duration of the reaction was 20 min with the
Materials 2019,12, 4075 3 of 17
following cooling in air. The properties of the samples prepared this way were described in a previous
paper [
28
]. The microstructure is composed of the two phases (NiTi austenite – cubic, Ti
2
Ni – cubic),
hardness 276 HV10, area fraction of the Ti
2
Ni phase 11.7%, and transformation temperatures: A
S
=
56
C, A
F
=86
C, and M
S
=21
C. The SHS product was milled in a vibratory cylinder mill VM4 (OPS
rerov, Pˇrerov, Czech Republic) in atmosphere with a duration of 7 min and the powder fraction with
a particle size <355
µ
m was selected by sieving using Fritsch Analysette 3 device (FRITSCH GmbH,
Germany). This pre-alloyed NiTi46 powder was consolidated by using the SPS method (FCT Systeme
HP D 10, Frankenblicke, Germany) at three various temperatures (900, 1000, and 1100
C) under the
pressure of 50 MPa with a holding time of 10 min. The high heating rate was chosen as 300
C/min at the
beginning and the last 100
C with the heating rate of 100
C/min (for example, sintering temperature
of 1000
C: applied heating rate of 300
C/min up to 900
C, between the temperatures 900 and 1000
C
and the applied heating rate of 100
C/min). The conditions of SPS consolidation as the temperature
regime, compaction force, height reduction, and current flow are shown in Figure 1.
Materials 2019, 12, x FOR PEER REVIEW 3 of 17
way were described in a previous paper [28]. The microstructure is composed of the two phases (NiTi
austenite – cubic, Ti2Ni – cubic), hardness 276 HV10, area fraction of the Ti2Ni phase 11.7%, and
transformation temperatures: AS = 56 °C, AF = 86 °C, and MS = 21 °C. The SHS product was milled in a
vibratory cylinder mill VM4 (OPS Přerov, Přerov, Czech Republic) in atmosphere with a duration of 7
minutes and the powder fraction with a particle size <355 µm was selected by sieving using Fritsch
Analysette 3 device (FRITSCH GmbH, Germany). This pre-alloyed NiTi46 powder was consolidated
by using the SPS method (FCT Systeme HP D 10, Frankenblicke, Germany) at three various
temperatures (900, 1000, and 1100 °C) under the pressure of 50 MPa with a holding time of 10 minutes.
The high heating rate was chosen as 300 °C/min at the beginning and the last 100 °C with the heating
rate of 100 °C/min (for example, sintering temperature of 1000 °C: applied heating rate of 300 °C/
min up to 900 °C, between the temperatures 900
Figure 1. SPS parameters during consolidation at the temperature of 900 °C.
The heat treatment in the temperature range from 400 °C to 900 °C was applied to the SPS-ed
samples. The duration of heat treatment was 60 minutes. Two variants of cooling were used including
a high cooling rate with quenching in water and a slow cooling rate, which was provided by cooling
in the closed furnace (average cooling rate approximately 2.5 °C/min between 700 °C and 300 °C).
The metallographic samples were prepared by grinding and polishing and the microstructure
was revealed by etching in Kroll’s reagent (5 mL HNO3, 10 mL HF, and 85 mL H2O). The
microstructure was observed using scanning electron microscopes (SEM) equipped with the EDS
(Energy Dispersive Spectroscopy) analyzers for identification of the chemical composition of the
individual phases: VEGA 3 LMU (TESCAN, Brno, Czech Republic) equipped with the OXFORD
Instruments X-max 20 mm2 SDD EDS analyzer (Oxford Instruments, HighWycombe, UK) and JEOL
IT 500 HR 500 (JEOL, Tokyo, Japan). The phase composition was analyzed by the X-ray diffraction
analysis (XRD) using a X’Pert Pro (PANalytical, Almelo, Netherlands) X-ray diffractometer with
CuKα radiation and a LynxEye XE detector (PANalytical, Almelo, The Netherlands). The mechanical
properties of the samples were evaluated by measuring Vickers hardness with a load of 10 kg and
compression tests (LabTest 5.250SP1-VM universal loading machine Labortech, Opava, Czech
Republic) with a strain rate of 0.3 s1 on samples measuring 3.3 mm × 3.3 mm × 5 mm. Compression
tests were conducted in both the direction (longitudinal and perpendicular) to the direction of SPS.
A longitudinal direction is parallel to compressive force during the SPS process. Differential scanning
Figure 1. SPS parameters during consolidation at the temperature of 900 C.
The heat treatment in the temperature range from 400
C to 900
C was applied to the SPS-ed
samples. The duration of heat treatment was 60 min. Two variants of cooling were used including a
high cooling rate with quenching in water and a slow cooling rate, which was provided by cooling in
the closed furnace (average cooling rate approximately 2.5 C/min between 700 C and 300 C).
The metallographic samples were prepared by grinding and polishing and the microstructure was
revealed by etching in Kroll’s reagent (5 mL HNO
3
, 10 mL HF, and 85 mL H
2
O). The microstructure
was observed using scanning electron microscopes (SEM) equipped with the EDS (Energy Dispersive
Spectroscopy) analyzers for identification of the chemical composition of the individual phases: VEGA
3 LMU (TESCAN, Brno, Czech Republic) equipped with the OXFORD Instruments X-max 20 mm
2
SDD EDS analyzer (Oxford Instruments, HighWycombe, UK) and JEOL IT 500 HR 500 (JEOL, Tokyo,
Japan). The phase composition was analyzed by the X-ray diraction analysis (XRD) using a X’Pert
Pro (PANalytical, Almelo, The Netherlands) X-ray diractometer with CuK
α
radiation and a LynxEye
XE detector (PANalytical, Almelo, The Netherlands). The mechanical properties of the samples
were evaluated by measuring Vickers hardness with a load of 10 kg and compression tests (LabTest
5.250SP1-VM universal loading machine Labortech, Opava, Czech Republic) with a strain rate of
Materials 2019,12, 4075 4 of 17
0.3 s
1
on samples measuring 3.3 mm
×
3.3 mm
×
5 mm. Compression tests were conducted in both
the direction (longitudinal and perpendicular) to the direction of SPS. A longitudinal direction is
parallel to compressive force during the SPS process. Dierential scanning calorimetry (DSC) analysis
of the prepared alloys was performed using Setaram DSC 131 (Setaram, Caluire, France) to determine
the transformation temperatures in products. Measurements for determining temperatures austenite
start (A
S
) and austenite finish (A
F
) were carried out in the temperature range of
20
C to 200
C at a
heating rate of 10
C/min and cooling from 200
C to
5
C for detecting the martensite start (M
S
) and
martensite finish (MF) temperatures.
The samples compacted by SPS at 900
C and with following heat treatments at 600 and 700
C
for 1 h and slow cooling in the closed furnace were also investigated using transmission electron
microscopy (JEOL JEM 2200FS, JEOL, Tokyo, Japan, accelerating voltage of 200 kV). Standard 3 mm
samples prepared by a slow-speed diamond blade cutting were mechanically dimpled and ion polished
in a Gatan PIPS 691 device (Pleasanton, CA, USA).
3. Results and Discussion
3.1. Microstructure, Phase Composition, and Phase Transformation
First, the influence of used sintering temperature on the quality of sintering individual particles
was investigated by a porosity measurement. Since it is visible in Figure 2and acquired by a light
microscope, there are dierences between samples compacted at various temperatures. The direction
of observation plays an important role. The non-deformed grains were observed on the perpendicular
cut (perpendicular to the direction of compression) whereas the elongated shape of grains after loading
during SPS remained in the microstructure on the longitudinal cut (parallel to SPS compression). The
highest values of porosity were determined at the samples sintered at 900
C. The porosity was high at
the perpendicular level and also at the longitudinal cut. The SPS temperature of 1000
C was sucient
to the reduction of porosity in comparison to the temperature of 900
C. Mainly in the longitudinal
direction, the value of porosity decreased rapidly to a similar value, which was measured after sintering
at 1100
C. The values of porosity are compared in Table 1and, based on the porosity measurement, it is
clear that the higher temperature of the SPS process leads to superior sintering of individual particles.
Materials 2019, 12, x FOR PEER REVIEW 4 of 17
calorimetry (DSC) analysis of the prepared alloys was performed using Setaram DSC 131 (Setaram,
Caluire, France) to determine the transformation temperatures in products. Measurements for
determining temperatures austenite start (AS) and austenite finish (AF) were carried out in the
temperature range of 20 °C to 200 °C at a heating rate of 10 °C/min and cooling from 200 °C to 5 °C
for detecting the martensite start (MS) and martensite finish (MF) temperatures.
The samples compacted by SPS at 900 °C and with following heat treatments at 600 and 700 °C
for 1 hour and slow cooling in the closed furnace were also investigated using transmission electron
microscopy (JEOL JEM 2200FS, JEOL, Tokyo, Japan, accelerating voltage of 200 kV). Standard 3 mm
samples prepared by a slow-speed diamond blade cutting were mechanically dimpled and ion
polished in a Gatan PIPS 691 device (Pleasanton, CA, USA).
3. Results and Discussion
3.1. Microstructure, Phase Composition, and Phase Transformation
First, the influence of used sintering temperature on the quality of sintering individual particles
was investigated by a porosity measurement. Since it is visible in Figure 2 and acquired by a light
microscope, there are differences between samples compacted at various temperatures. The direction
of observation plays an important role. The non-deformed grains were observed on the perpendicular
cut (perpendicular to the direction of compression) whereas the elongated shape of grains after
loading during SPS remained in the microstructure on the longitudinal cut (parallel to SPS
compression). The highest values of porosity were determined at the samples sintered at 900 °C. The
porosity was high at the perpendicular level and also at the longitudinal cut. The SPS temperature of
1000 °C was sufficient to the reduction of porosity in comparison to the temperature of 900 °C. Mainly
in the longitudinal direction, the value of porosity decreased rapidly to a similar value, which was
measured after sintering at 1100 °C. The values of porosity are compared in Table 1 and, based on the
porosity measurement, it is clear that the higher temperature of the SPS process leads to superior
sintering of individual particles.
(a)
(b)
Figure 2. Cont.
Materials 2019,12, 4075 5 of 17
Materials 2019, 12, x FOR PEER REVIEW 5 of 17
(c)
(d)
(e)
(f)
Figure 2. Microstructure and porosity of the NiTi46 alloy SPS consolidated at various
temperatures: (a) 900 °C-perpendicular, (b) 900 °C-longitudinal, (c) 1000 °C-perpendicular, (d) 1000
°C-longitudinal, (e) 1100 °C-perpendicular, and (f) 1100 °C-longitudinal direction.
Table 1. Influence of SPS temperature on porosity of the sample and area fraction of the Ti2Ni phase
in the microstructure.
Parameter Direction Spark Plasma Sintering Temperature
900 °C 1000 °C 1100 °C
Porosity (%) Perpendicular 1.6 ± 0.09 0.7 ± 0.16 ˂ 0.1
Longitudinal 1.4 ± 0.09 0.1 ± 0.06 ˂ 0.1
Area fraction of the Ti2Ni phase (%) 13.8 ± 2.41 17.0 ± 2.22 11.9 ± 1.22
The SPS temperature influences the phase composition and also mechanical properties. The
effect of sintering temperature was investigated in the previous paper [29]. The area fraction of the
undesirable Ti2Ni and Ni3Ti phases increased with an increasing sintering temperature. The high
amount of the Ti2Ni phase was formed along the boundaries of the sintered particles. It is necessary
to point out that, in the previous paper, pulse current flow through the sample (another SPS device)
was used while, in this paper, the regime of direct current flow through the sample was applied and
it is the reason for different results. In this case, the lower amounts of the Ti2Ni phase were measured
generally and a growing trend with SPS temperature was not observed. The area fraction of the Ti2Ni
phase increased slightly by SPS consolidation at 900 °C and 1000 °C, but, after SPS sintering at 1100
Figure 2.
Microstructure and porosity of the NiTi46 alloy SPS consolidated at various temperatures:
(
a
) 900
C-perpendicular, (
b
) 900
C-longitudinal, (
c
) 1000
C-perpendicular, (
d
) 1000
C-longitudinal,
(e) 1100 C-perpendicular, and (f) 1100 C-longitudinal direction.
Table 1.
Influence of SPS temperature on porosity of the sample and area fraction of the Ti
2
Ni phase in
the microstructure.
Parameter Direction Spark Plasma Sintering Temperature
900 C 1000 C 1100 C
Porosity (%) Perpendicular 1.6 ±0.09 0.7 ±0.16 <0.1
Longitudinal 1.4 ±0.09 0.1 ±0.06 <0.1
Area fraction of the Ti2Ni phase (%) 13.8 ±2.41 17.0 ±2.22 11.9 ±1.22
The SPS temperature influences the phase composition and also mechanical properties. The
eect of sintering temperature was investigated in the previous paper [
29
]. The area fraction of the
undesirable Ti
2
Ni and Ni
3
Ti phases increased with an increasing sintering temperature. The high
amount of the Ti
2
Ni phase was formed along the boundaries of the sintered particles. It is necessary
to point out that, in the previous paper, pulse current flow through the sample (another SPS device)
was used while, in this paper, the regime of direct current flow through the sample was applied and it
is the reason for dierent results. In this case, the lower amounts of the Ti
2
Ni phase were measured
generally and a growing trend with SPS temperature was not observed. The area fraction of the Ti
2
Ni
phase increased slightly by SPS consolidation at 900
C and 1000
C, but, after SPS sintering at 1100
C,
Materials 2019,12, 4075 6 of 17
the amount of the Ti
2
Ni phase was reduced to approximately 12%, which means a comparable value to
result in samples prepared by the SHS method [29] (see Table 1).
At all SPS temperatures, the SEM observation was performed. Improving the fusion of grain
boundaries was confirmed with increasing sintering temperature. The NiTi phase matrix with the
fine Ni-rich precipitates in all samples was commonly found within the Ti
2
Ni and Ni
3
Ti phases.
The microstructures after SPS are shown in Figure 3. In Table 2, there are summarized chemical
compositions of individual areas observed by SEM. A good agreement in chemical compositions of the
NiTi and Ti
2
Ni phases to the binary Ni-Ti phase diagram was found out. The chemical composition of
the area labelled 3 is close to the Ni
3
Ti phase. Chemical composition of areas 7 and 10 is approaching
the chemical composition of the Ni4Ti3phase.
A more detailed observation of the microstructure was performed using a transmission electron
microscope (TEM). The main goal of this experiment is based on investigating the fine needle-like
particles in the NiTi matrix. Figure 4shows TEM micrographs of sample SPS-ed at the temperature of
900
C. The NiTi and Ti
2
Ni phase were observed commonly with the Ni
4
Ti
3
phase (determined by
electron diraction) in the NiTi matrix.
Figure 3. Cont.
Materials 2019,12, 4075 7 of 17
Materials 2019, 12, x FOR PEER REVIEW 7 of 17
(e)
(f)
Figure 3. SEM micrographs of SPS-ed samples at various temperatures: (a) 900 °C, (b) 900 °C-detail
of Ni-rich precipitates, (c) 1000 °C, (d) 1000 °C-detail of Ni-rich precipitates, (e) 1100 °C, and (f)
1100 °C-detail of Ni-rich precipitates and formed the Ni3Ti phase.
Table 2. Chemical composition of individual areas measured by EDS analysis.
Area Ni (wt%) Ti (wt%)
1 54.4 45.6
2 38.0 62.0
3 72.4 27.6
4 68.6 31.4
5 55.5 44.5
6 38.0 62.0
7 58.7 41.3
8 55.2 44.8
9 37.6 62.4
10 59.1 40.9
(a)
(b)
Figure 3.
SEM micrographs of SPS-ed samples at various temperatures: (
a
) 900
C, (
b
) 900
C-detail
of Ni-rich precipitates, (
c
) 1000
C, (
d
) 1000
C-detail of Ni-rich precipitates, (
e
) 1100
C, and
(f) 1100 C-detail of Ni-rich precipitates and formed the Ni3Ti phase.
Table 2. Chemical composition of individual areas measured by EDS analysis.
Area Ni (wt.%) Ti (wt.%)
1 54.4 45.6
2 38.0 62.0
3 72.4 27.6
4 68.6 31.4
5 55.5 44.5
6 38.0 62.0
7 58.7 41.3
8 55.2 44.8
9 37.6 62.4
10 59.1 40.9
Materials 2019, 12, x FOR PEER REVIEW 7 of 17
(e)
(f)
Figure 3. SEM micrographs of SPS-ed samples at various temperatures: (a) 900 °C, (b) 900 °C-detail
of Ni-rich precipitates, (c) 1000 °C, (d) 1000 °C-detail of Ni-rich precipitates, (e) 1100 °C, and (f)
1100 °C-detail of Ni-rich precipitates and formed the Ni3Ti phase.
Table 2. Chemical composition of individual areas measured by EDS analysis.
Area Ni (wt%) Ti (wt%)
1 54.4 45.6
2 38.0 62.0
3 72.4 27.6
4 68.6 31.4
5 55.5 44.5
6 38.0 62.0
7 58.7 41.3
8 55.2 44.8
9 37.6 62.4
10 59.1 40.9
(a)
(b)
Figure 4. Cont.
Materials 2019,12, 4075 8 of 17
Materials 2019, 12, x FOR PEER REVIEW 8 of 17
(c)
Figure 4. Ni4Ti3 phase in the matrix: (a) bright-field micrograph, g = 10-1, close to the [4,1,4] zone
axis, (b) corresponding dark-field micrograph using Ni4Ti3 spot marked by a circle in the
diffraction pattern in the inset, and (c) bright filed micrograph showing dark Ni4Ti3 particles in the
matrix adjacent to a Ti2Ni particle.
The effect of heat treatment on the microstructure and the phase composition was investigated
at samples SPS-ed at the temperature of 900 °C. The phase compositions of SPS-ed samples were
verified by XRD analysis. The diffraction lines were very similar through all SPS consolidation
temperatures. The phase composition of SPS-ed samples consists of the NiTi phases (Cubic, Pm-3m),
Ti2Ni phase (Cubic, Fd-3 m), Ni3Ti phase (Hexagonal, P63/mmc), and the Ni4Ti3 phase (Rhomboedral,
R-3). XRD patterns are displayed. After heat treatment in the temperature range of 600–900 °C with
a 1-hour duration, there were no observed changes in the phase compositions of the samples (see
Figure 5). This fact can seem to be strange in comparison with other studies. However, it is necessary
to consider the initial state in individual studies (mostly after homogenization annealing, e.g.,
[17,27]). The initial state of these samples is after SPS consolidation at 900, 1000, and 1100 °C. The
very high heating rate (approximately 300 °C/min) and a very high cooling rate (as shown in Figure
1) was applied during the SPS process, whereas the initial state of the Ni-Ti alloys is after
homogenization or solution annealing for tens of minutes or several hours at temperatures at around
1000 °C [17,27,30]. Moreover, the transformations in the microstructure and phase composition occur
in a bigger extent after heat treatment with longer duration since it is clear in this study [17]. The
crystallite sizes of the NiTi phase after spark plasma sintering, which was determined by the means
of the Sherrer’s formula, range from 20 to 47 nm. The crystallite sizes increased after heat treatment
and the highest values of the crystallite sizes were determined at samples heat-treated at a
temperature of 700 °C. The significant increase of the crystallite sizes of the NiTi phase is related
likely to the recrystallization process, grain growth, or the order-disorder transformation around the
temperature of 600–700 °C reported in References [31,32]. The values of the crystallite sizes of the
NiTi phase depending on the regime of the heat treatment are stated in Table 3.
Figure 4.
Ni
4
Ti
3
phase in the matrix: (
a
) bright-field micrograph, g =10-1, close to the [4,1,4] zone axis,
(
b
) corresponding dark-field micrograph using Ni
4
Ti
3
spot marked by a circle in the diraction pattern
in the inset, and (
c
) bright filed micrograph showing dark Ni
4
Ti
3
particles in the matrix adjacent to a
Ti2Ni particle.
The eect of heat treatment on the microstructure and the phase composition was investigated at
samples SPS-ed at the temperature of 900
C. The phase compositions of SPS-ed samples were verified
by XRD analysis. The diraction lines were very similar through all SPS consolidation temperatures.
The phase composition of SPS-ed samples consists of the NiTi phases (Cubic, Pm-3m), Ti
2
Ni phase
(Cubic, Fd-3 m), Ni
3
Ti phase (Hexagonal, P63/mmc), and the Ni
4
Ti
3
phase (Rhomboedral, R-3). XRD
patterns are displayed. After heat treatment in the temperature range of 600–900
C with a 1-h duration,
there were no observed changes in the phase compositions of the samples (see Figure 5). This fact
can seem to be strange in comparison with other studies. However, it is necessary to consider the
initial state in individual studies (mostly after homogenization annealing, e.g., [
17
,
27
]). The initial
state of these samples is after SPS consolidation at 900, 1000, and 1100
C. The very high heating
rate (approximately 300
C/min) and a very high cooling rate (as shown in Figure 1) was applied
during the SPS process, whereas the initial state of the Ni-Ti alloys is after homogenization or solution
annealing for tens of minutes or several hours at temperatures at around 1000
C [
17
,
27
,
30
]. Moreover,
the transformations in the microstructure and phase composition occur in a bigger extent after heat
treatment with longer duration since it is clear in this study [
17
]. The crystallite sizes of the NiTi phase
after spark plasma sintering, which was determined by the means of the Sherrer’s formula, range from
20 to 47 nm. The crystallite sizes increased after heat treatment and the highest values of the crystallite
sizes were determined at samples heat-treated at a temperature of 700
C. The significant increase of
the crystallite sizes of the NiTi phase is related likely to the recrystallization process, grain growth, or
the order-disorder transformation around the temperature of 600–700
C reported in References [
31
,
32
].
The values of the crystallite sizes of the NiTi phase depending on the regime of the heat treatment are
stated in Table 3.
Materials 2019,12, 4075 9 of 17
Materials 2019, 12, x FOR PEER REVIEW 9 of 17
(a)
(b)
(c)
Figure 5. XRD patterns of the NiTi46 alloys: (a) spark plasma sintered at 900, 1000, and 1100 °C, (b)
spark plasma sintered at 900 °C, heat-treated at 600–900 °C and slowly cooled, (c) spark plasma
sintered at 900 °C, heat-treated at 600–900 °C, and cooled in water.
Table 3. Crystallite sizes of the NiTi phase depending on the regime of heat treatment.
Figure 5.
XRD patterns of the NiTi46 alloys: (
a
) spark plasma sintered at 900, 1000, and 1100
C,
(
b
) spark plasma sintered at 900
C, heat-treated at 600–900
C and slowly cooled, (
c
) spark plasma
sintered at 900 C, heat-treated at 600–900 C, and cooled in water.
Materials 2019,12, 4075 10 of 17
Table 3. Crystallite sizes of the NiTi phase depending on the regime of heat treatment.
Sample Crystallite Size (nm)
SPS 900 C 47
SPS 1000 C 34
SPS 1100 C 20
SPS 900 C-HT 600 C-furnace 25
SPS 900 C-HT 700 C-furnace 122
SPS 900 C-HT 900 C-furnace 28
SPS 900 C-HT 600 C-water 84
SPS 900 C-HT 700 C-water 129
SPS 900 C-HT 900 C-water 57
From the point of view of the microstructure, the particles of the Ti
2
Ni, Ni
3
Ti phases and Ni
4
Ti
3
needles in the NiTi matrix were observed in Figure 6. A higher amount of the Ni
3
Ti phase was formed
in the sample heat-treated at 700
C and cooled in water than in the samples with slow cooling in
the closed furnace after heat treatment. The area fraction of the Ti
2
Ni phase after heat treatment was
almost invariable with values between 13% to 16%. Detailed observation of the Ni
4
Ti
3
and other
phases after heat treatment at 600 C and 700 C was carried out by TEM again (see Figure 7).
Materials 2019, 12, x FOR PEER REVIEW 10 of 17
Sample Crystallite Size (nm)
SPS 900 °C 47
SPS 1000 °C 34
SPS 1100 °C 20
SPS 900 °C-HT 600 °C-furnace 25
SPS 900 °C-HT 700 °C-furnace 122
SPS 900 °C-HT 900 °C-furnace 28
SPS 900 °C-HT 600 °C-water 84
SPS 900 °C-HT 700 °C-water 129
SPS 900 °C-HT 900 °C-water 57
From the point of view of the microstructure, the particles of the Ti2Ni, Ni3Ti phases and Ni4Ti3
needles in the NiTi matrix were observed in Figure 6. A higher amount of the Ni3Ti phase was formed
in the sample heat-treated at 700 °C and cooled in water than in the samples with slow cooling in the
closed furnace after heat treatment. The area fraction of the Ti2Ni phase after heat treatment
was almost invariable with values betw
(a) (b)
(c) (d)
Figure 6.
SEM microstructures of samples SPS-ed at 900
C and heat-treated: (
a
,
c
) at 700
C followed
with slow cooling in the closed furnace, (
b
) at 700
C and cooled in water, and (
d
) at 600
C followed
with slow cooling in the closed furnace.
Materials 2019,12, 4075 11 of 17
Materials 2019, 12, x FOR PEER REVIEW 11 of 17
Figure 6. SEM microstructures of samples SPS-ed at 900 °C and heat-treated: (a,c) at 700 °C
followed with slow cooling in the closed furnace, (b) at 700 °C and cooled in water, and (d) at 600
°C followed with slow cooling in the closed furnace.
(a)
(b)
(c)
(d)
Figure 7. TEM observation of samples heat-treated at 600 °C (a,b) and at 700 °C (c,d) and slow
cooled: (a) Ni4Ti3 precipitates with a typical lenticular cross-section, (b) corresponding diffraction
pattern in the zone axis [2,-1-1,3] of the precipitate, which is coincident with the zone axis [1,0,2] of
the B2 NiTi matrix, (c) Ni4Ti3 particles in the NiTi matrix adjacent to very fine grains of the Ti2Ni
phase, (d) Ni4Ti3 particle adjacent to a coarser grain of the Ti2Ni phase. The inset shows a diffraction
pattern of the Ti2Ni phase in the [1,1,2] zone axis.
The phase transformation in the NiTi phase between austenite and the martensite structure was
studied using differential scanning calorimetry. The straight lines were obtained for the samples as-
SPS sintered at all temperatures and any phase transformation occurs in these samples. The change
in the phase transformation behaviour was brought by heat treatment of the samples. When the
samples were heat-treated, the temperature of heat treatment and way of cooling are important. Fast
cooling in water did not cause the recovery of the phase transformation. Therefore, the microstructure
observation and mechanical property investigation are focused on the samples cooled slowly in the
Figure 7.
TEM observation of samples heat-treated at 600
C (
a
,
b
) and at 700
C (
c
,
d
) and slow cooled:
(
a
) Ni
4
Ti
3
precipitates with a typical lenticular cross-section, (
b
) corresponding diraction pattern in
the zone axis [2,-1-1,3] of the precipitate, which is coincident with the zone axis [1,0,2] of the B2 NiTi
matrix, (
c
) Ni
4
Ti
3
particles in the NiTi matrix adjacent to very fine grains of the Ti
2
Ni phase, (
d
) Ni
4
Ti
3
particle adjacent to a coarser grain of the Ti
2
Ni phase. The inset shows a diraction pattern of the Ti
2
Ni
phase in the [1,1,2] zone axis.
The phase transformation in the NiTi phase between austenite and the martensite structure was
studied using dierential scanning calorimetry. The straight lines were obtained for the samples as-SPS
sintered at all temperatures and any phase transformation occurs in these samples. The change in the
phase transformation behaviour was brought by heat treatment of the samples. When the samples were
heat-treated, the temperature of heat treatment and way of cooling are important. Fast cooling in water
did not cause the recovery of the phase transformation. Therefore, the microstructure observation and
mechanical property investigation are focused on the samples cooled slowly in the closed furnace. The
peaks on DSC curves were formed only at samples after heat treatment at temperatures of 600–700
C,
which were slow cooled in the closed furnace. The heating and cooling DSC curves of samples SPS-ed
at 900 C are shown in Figure 8.
Materials 2019,12, 4075 12 of 17
Materials 2019, 12, x FOR PEER REVIEW 12 of 17
closed furnace. The peaks on DSC curves were formed only at samples after heat treatment at
temperatures of 600–700 °C, which were slow cooled in the closed furnace. The heating and cooling
DSC curves of samples SPS-ed at 900 °C are shown in Figure 8.
(a)
(b)
(c)
(d)
Figure 8. DSC heating and cooling curves: (a) Heating curves of samples SPS-ed at 900 °C and heat-
treated between 500 and 700 °C. (b) Cooling curves of samples SPS-ed at 900 °C and heat-treated
between 500 and 700 °C. (c) Heating curves of samples SPS-ed at 1000 and 1100 °C and SPS-ed at
1000 and 1100 °C with heat treatment at 600 °C. (d) Cooling curves of samples SPS-ed at 1000 and
1100 °C SPS-ed at 1000 and 1100 °C with heat treatment at 600 °C.
3.2. Mechanical Properties
The increase of hardness with the increasing temperature of SPS is similar to the previous results
[29]. However, an increase of hardness was assigned to the rising value of the Ti2Ni phase. It is
contrary to the current study and to the decrease of the Ti2Ni phase amount after SPS at 1100 °C.
There are two possible reasons. Firstly, the hardness increases with the quality of powder sintering
at higher SPS temperature, which is visible in Figure 3. Secondly, the precipitation process of Ni-rich
phases occurs during the SPS process while heat treatment at the temperatures of 900–1000 °C lead
to improved hardness [17]. Values of hardness as SPS-ed samples are stated in Table 4.
Figure 8.
DSC heating and cooling curves: (
a
) Heating curves of samples SPS-ed at 900
C and
heat-treated between 500 and 700
C. (
b
) Cooling curves of samples SPS-ed at 900
C and heat-treated
between 500 and 700
C. (
c
) Heating curves of samples SPS-ed at 1000 and 1100
C and SPS-ed at 1000
and 1100
C with heat treatment at 600
C. (
d
) Cooling curves of samples SPS-ed at 1000 and 1100
C
SPS-ed at 1000 and 1100 C with heat treatment at 600 C.
3.2. Mechanical Properties
The increase of hardness with the increasing temperature of SPS is similar to the previous
results [
29
]. However, an increase of hardness was assigned to the rising value of the Ti
2
Ni phase. It
is contrary to the current study and to the decrease of the Ti
2
Ni phase amount after SPS at 1100
C.
There are two possible reasons. Firstly, the hardness increases with the quality of powder sintering at
higher SPS temperature, which is visible in Figure 3. Secondly, the precipitation process of Ni-rich
phases occurs during the SPS process while heat treatment at the temperatures of 900–1000
C lead to
improved hardness [17]. Values of hardness as SPS-ed samples are stated in Table 4.
Materials 2019,12, 4075 13 of 17
Table 4.
Summary of hardness and compressive stress-strain test of samples after spark plasma
sintering at various temperatures. Longitudinal direction is parallel to the direction of compressive
force by the SPS process.
SPS Temperature 900 C 1000 C 1100 C
Hardness (HV 10)/Std. dev. (±) 562/25 596/20 624/23
Longitudinal UCS (MPa) 1903 2116 2315
Agt (%) 8.7 8.7 8.7
Perpendicular UCS (MPa) 1953 2212 2243
Agt (%) 7.4 9.4 8.6
The stress-strain behavior was analyzed in tandem with hardness and the same influence
(increasing values of UCS) with an increasing temperature of the SPS process was observed. As
visible in Figure 9, the samples after SPS reach high values of ultimate compressive strength (UCS)
1900–2300 MPa, but there are not the areas of plastic deformation on stress-strain curves and the samples
fail in a brittle manner at a maximum load perpendicularly as well as in the longitudinal direction.
The values of elongation at maximum force (Agt) were between 7.4–9.4% for all SPS temperatures and
both directions. In the case of the compression test, the lower porosity of samples SPS-ed at higher
temperatures is likely connected with increasing values of UCS.
After heat treatment with cooling in the closed furnace, the decrease of hardness was found and
the lowest value was measured after heat treatment at the temperature of 700
C. Currently, with the
decrease of hardness, increased ductility and plasticity of samples was measured by a compressive
test (Table 5). The evolution of hardness depending on the temperature of heat treatment is shown in
Figure 9for samples SPS sintered at 900
C. In case of samples SPS sintered, a 1000
C and 1100
C,
hardness after heat treatment at 600
C dropped to values of 504 HV 10 and 509 HV 10. By the
compressive test after heat treatment, the dierence between annealing temperatures of 600
C and
700
C was observed. In case of annealing temperature of 600
C, the similar values of Agt were
measured at longitudinal and perpendicular directions between 9.5% and 10.5%, whereas, at an
annealing temperature of 700
C, the increase of Agt was found and the dierence between longitudinal
(16.3–19.1%) and perpendicular (9.3–13.6%) direction grew up. The UCS of the samples SPS-ed at
1000
C and 1100
C increased after heat treatment about 100 MPa in comparison with the state after
SPS. Generally, the values of UCS and Agt increased with an increasing temperature of the SPS process.
Materials 2019, 12, x FOR PEER REVIEW 13 of 17
Table 4. Summary of hardness and compressive stress-strain test of samples after spark plasma
sintering at various temperatures. Longitudinal direction is parallel to the direction of compressive
force by the SPS process.
SPS Temperature 900 °C 1000 °C 1100 °C
Hardness (HV 10)/Std. dev. (±) 562/25 596/20 624/23
Longitudinal UCS (MPa) 1903 2116 2315
Agt (%) 8.7 8.7 8.7
Perpendicular UCS (MPa) 1953 2212 2243
Agt (%) 7.4 9.4 8.6
The stress-strain behavior was analyzed in tandem with hardness and the same influence
(increasing values of UCS) with an increasing temperature of the SPS process was observed. As
visible in Figure 9, the samples after SPS reach high values of ultimate compressive strength (UCS)
1900 – 2300 MPa, but there are not the areas of plastic deformation on stress-strain curves and the
samples fail in a brittle manner at a maximum load perpendicularly as well as in the longitudinal
direction. The values of elongation at maximum force (Agt) were between 7.4–9.4% for all SPS
temperatures and both directions. In the case of the compression test, the lower porosity of samples
SPS-ed at higher temperatures is likely connected with increasing values of UCS.
After heat treatment with cooling in the closed furnace, the decrease of hardness was found and
the lowest value was measured after heat treatment at the temperature of 700 °C. Currently, with the
decrease of hardness, increased ductility and plasticity of samples was measured by a compressive
test (Table 5). The evolution of hardness depending on the temperature of heat treatment is shown in
Figure 9 for samples SPS sintered at 900 °C. In case of samples SPS sintered, a 1000 °C and 1100 °C,
hardness after heat treatment at 600 °C dropped to values of 504 HV 10 and 509 HV 10. By the
compressive test after heat treatment, the difference between annealing temperatures of 600 °C and
700 °C was observed. In case of annealing temperature of 600 °C, the similar values of Agt were
measured at longitudinal and perpendicular directions between 9.5% and 10.5%, whereas, at an
annealing temperature of 700 °C, the increase of Agt was found and the difference between
longitudinal (16.3–19.1%) and perpendicular (9.3–13.6%) direction grew up. The UCS of the samples
SPS-ed at 1000 °C and 1100 °C increased after heat treatment about 100 MPa in comparison with the
state after SPS. Generally, the values of UCS and Agt increased with an increasing temperature of the
SPS process.
(a)
(b)
Figure 9. Cont.
Materials 2019,12, 4075 14 of 17
Materials 2019, 12, x FOR PEER REVIEW 14 of 17
(c)
(d)
(e)
Figure 9. Compressive stress-strain curves and evolution of hardness: (a) Stress-strain curves of
samples SPS-ed at 900 °C and heat-treated-perpendicular direction. (b) Stress-strain curves of
samples SPS-ed at 900 °C and heat-treated-longitudinal direction. Stress-strain curves of samples
SPS-ed at 1000 °C, 1100 °C, and heat-treated-perpendicular (c) and longitudinal (d) direction. (e)
Evolution of hardness during heat treatment of sample SPS-ed at 900 °C.
Table 5. Summary of mechanical properties of heat-treated samples consolidated by the SPS
process.
Sample Heat treatment regime Hardness (HV 10)/Std.
dev. (±) Direction UCS
(MPa)
Agt
(%)
SPS 900 °C 600 °C-furnace 423/21 Perpendicular 2163 10.5
Longitudinal 2089 11.6
SPS 900 °C 600 °C-water 531/14 Perpendicular 1508 7.6
Longitudinal 1994 9.9
SPS 900 °C 700 °C-furnace 388/20 Perpendicular 1430 9.3
Figure 9.
Compressive stress-strain curves and evolution of hardness: (
a
) Stress-strain curves of
samples SPS-ed at 900
C and heat-treated-perpendicular direction. (
b
) Stress-strain curves of samples
SPS-ed at 900
C and heat-treated-longitudinal direction. Stress-strain curves of samples SPS-ed at
1000
C, 1100
C, and heat-treated-perpendicular (
c
) and longitudinal (
d
) direction. (
e
) Evolution of
hardness during heat treatment of sample SPS-ed at 900 C.
Materials 2019,12, 4075 15 of 17
Table 5. Summary of mechanical properties of heat-treated samples consolidated by the SPS process.
Sample Heat Treatment
Regime
Hardness (HV 10)/Std.
dev. (±)Direction UCS (MPa) Agt (%)
SPS 900 C 600 C-furnace 423/21 Perpendicular 2163 10.5
Longitudinal 2089 11.6
SPS 900 C 600 C-water 531/14 Perpendicular 1508 7.6
Longitudinal 1994 9.9
SPS 900 C 700 C-furnace 388/20 Perpendicular 1430 9.3
Longitudinal 1907 16.3
SPS 900 C 700 C-water 450/19 Perpendicular 1566 9.3
Longitudinal 1878 13.7
SPS 900 C 900 C-furnace 429/22 Perpendicular 1957 11.1
SPS 900 C 900 C-water 550/20 Perpendicular 1670 -
SPS 1000 C 600 C-furnace 504/21 Perpendicular 2089 9.5
Longitudinal 2235 9.2
SPS 1000 C 700 C-furnace 444/14 Perpendicular 2099 12.5
Longitudinal 2355 18.5
SPS 1100 C 600 C-furnace 509/22 Perpendicular 2295 10.0
Longitudinal 2488 10.6
SPS 1100 C 700 C-furnace 448/10 Perpendicular 2163 13.6
Longitudinal 2465 19.1
The explanation of the rapid decrease of hardness could be caused due to removing the deformation
strengthening from the milling of an SHS product. This type of decrease has to have the same scale for
all samples. The dependence of hardness on the temperature of heat treatment is undeniable. Thus,
the dierent processes and changes in microstructure must occur during heat treatment at various
temperatures. However, these changes were not observed in phase composition and microstructure
due to the short time of annealing during heat treatment. In this study [
17
], the longer time heat
treatment was applied and the following changes in the microstructure occurred. The high UCS,
hardness, and low ductility of samples were attributed to the precipitates of Ni
3
Ti
4
and Ni
3
Ti
2
phases,
while, when the microstructure contains the Ni
3
Ti phase or combination of the Ni
3
Ti and Ni
3
Ti
2
phases (corresponding to heat treatment between 600 and 800
C while a longer time must be applied
at a temperature of 600
C), high mechanical properties and good ductility were obtained. Both
studies [
17
,
24
] show similar evolution of the hardness values during heat treatment when compared to
our results.
4. Conclusions
The fabrication process composed of Self-propagating High-temperature Synthesis (SHS) and
Spark Plasma Sintering (SPS) was chosen to obtain a completely dense material. The highest temperature
of SPS (1100
C) led to the highest values of ultimate compressive strength without the formation of an
excessive number of undesirable phases. The following heat treatment is necessary to obtain a material
with a good combination of strength and ductility. The heat treatment leads to the disappearance of
the deformation strengthening coming from milling in the vibration mill and to recover the phase
transformation between the austenite and martensite structure of the NiTi phase, which was detected by
dierential scanning calorimetry (DSC). The heat treatment near the temperatures of 600–700
C with
slow cooling is recommended to obtain good ductility, strength, and probable shape memory properties.
Author Contributions:
P.S. and P.N. designed the experiment and evaluated the phase composition. P.S. and A.Š.
provided sample preparation by the SHS reaction, measurement of mechanical properties, and DSC analysis. F.P.
consolidated samples by the SPS method. P.S. and J.D. analyzed the microstructure by LM and SEM. M.K. and P.H.
observed the microstructure by SEM and TEM. P.S. wrote the paper. P.N. and M.K. reviewed and edited the paper.
Funding:
European Regional Development Fund (projects Pre-Application Research of Functionally Graduated
Materials by Additive Technologies, No. CZ.02.1.01/0.0/0.0/17_048/0007350 and Nanomaterials Centre for Advanced
Materials 2019,12, 4075 16 of 17
Applications, No. CZ.02.1.01/0.0/0.0/15_003/0000485) and specific university research (MSMT No. 21-SVV/2019)
supported the research.
Conflicts of Interest: The authors declare no conflict of interest.
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... A significant part of the Special Issue is devoted to the functional materials, which are used for their shape memory properties, magnetic behaviour, optical properties, self-healing beahaviour or biodegradability in the organism. Salvetr et al. studied the synthesis of a NiTi shape memory alloy by self-propagating high-temperature synthesis, milling and spark plasma sintering [7]. The alloy reached almost the theoretical density, a very high ultimate compressive strength of more than 2000 MPa, but it lost the transformation behaviour. ...
... The alloy reached almost the theoretical density, a very high ultimate compressive strength of more than 2000 MPa, but it lost the transformation behaviour. Further heat treatment at the temperature of 600-700 • C with slow cooling was needed to recover the transformation behaviour, which is crucial for the shape memory effect [7]. The magnetic materials were investigated by Skotnicova et al. [8]. ...
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Powder metallurgy is a group of advanced processes for the synthesis, processing, and shaping of various kinds of materials. Initially inspired by ceramics processing, the methodology comprising of the production of a powder and its transformation to a compact solid product has attracted great attention since the end of World War II. At present, there are many technologies for powder production (e.g., gas atomization of the melt, chemical reduction, milling, and mechanical alloying) and its consolidation (e.g., pressing and sintering, hot isostatic pressing, and spark plasma sintering). The most promising ones can achieve an ultra-fine or nano-grained structure of the powder, and preserve it during consolidation. Among these methods, mechanical alloying and spark plasma sintering play a key role. This Special Issue gives special focus to the advancement of mechanical alloying, spark plasma sintering and self-propagating high-temperature synthesis methods, as well as to the role of these processes in the development of new materials.
... However, Ni 4 Ti 3 phase is a metastable phase, which would decompose at higher temperatures and with prolonged exposure. In addition to that, Ni 4 Ti 3 has a typical lenticular nanoscale shape [39,40]. Therefore, it was assumed to be the start phase of the decomposition inside the NiTi phase, which needs further TEM observation. ...
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... Similarly, the binary NiTi SMA has an electron concentration of around 7. Figure 1a reveals the XRD pattern of the non-equiatomic binary NiTi alloys. The peaks were indexed by the literature [33][34][35][36]. All cases, excluding NT1, show austenite phase at room temperature, while NT1 entirely has a martensite crystal structure. ...
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... It is attributed to the effect of work hardening [8]; nevertheless, the effect of NiTi alloys pseudoelastic behavior is mostly responsible for good wear resistance [9,10]. In addition, NiTi possesses damping-sensitive properties that make it a good candidate for material robots and smart materials actuators in the areas of micro-electro-mechanical system (MEMS) devices [11,12]. Although the hardness of NiTi SMAs is much lower, better wear resistance has been reported as compared to those of conventional wear-resistant materials such as steel and alloy [13,14]. ...
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... However, one of the limitations lies in the area of manufacture of complex shapes by the PM route. There are various routes of PM methods such as Conventional sintering (CS), Self-propagating high-temperature synthesis (SHS), Metal injection molding (MIM), Hot isostatic pressing (HIP), Spark plasma sintering (SPS), and Plasma melting (PM), microwave sintering (MWS) methods [19][20][21][22]. Out of these sintering processes, the spark plasma sintering method stands out due to the short processing time of the reaction process. ...
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... The transformation interface propagation and the nucleation process may be altered using vacancies, dislocations, grain boundaries, solute chemical elements or precipitated particles [33,34]. For NiTi SMAs fabricated from powders by Spark Plasma Sintering (SPS), subsequent heat treatment at temperatures from 600 to 700 • C with slow cooling was found to be necessary to obtain good ductility [35]. ...
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In order to develop the application of the more cost-effective copper-based shape memory alloys (SMAs), rather than nickel–titanium as earthquake energy dissipaters, the influence of ausforming-induced plastic deformation on phase transformations, microstructure, super elasticity and mechanical properties of the shape memory alloy Cu-26Zn-4Al was examined. These specific SMA properties were targeted by applying appropriate parameters of the thermomechanical (the so-called ausforming) process: beta-phase homogenization at 800 °C for 20 min, one-step hot rolling at 800 °C and water quenching. The results showed significant microstructural changes, increased mechanical resistance and change in the phase transformation behavior. The SMA treated by ausforming retained the reversible austenitic–martensitic transformation ability, with the appearance of the super-elastic effect up to 6% of strain recovery. Although some strengthening occurred after hot rolling (an increase in true yield strength of 125 MPa was detected), all phase transformation temperatures were decreased. The smallest decrease was detected for the austenite finish temperature (32.8 °C) and the largest for the martensite finish temperature (42.0 °C), allowing both the expansion and the lowering of the temperature range of super elasticity, which is favorable for construction applications. It is concluded that it is possible to achieve an optimal combination of adequate strength and improved transformation behavior of Cu-Zn-Al alloy by applying the ausforming treatment.
... Optical microscopy ( Fig. 2) showed that the lower relative density HP samples (98.7%) were found to have more frequent pores than the higher density LP samples (99.5%). While no apparent martensite phases were observed in the HP samples, the LP samples contained a higher amount of needlelike martensite phases and a small amount of needle shaped precipitates probably Ni 4 Ti 3 (magnified area in Fig. 4e and g) which have higher nickel contents than other area and showed a similar morphology to previous studies [38,39]. ...
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This work is focused on the possibilities of preparing Ni-Ti46 wt pct alloy by powder metallurgy methods. The self-propagating high-temperature synthesis (SHS) and combination of SHS reaction, milling, and spark plasma sintering consolidation (SPS) are explored. The aim of this work is the development of preparation method with the lowest amount of undesirable phases (mainly Ti2Ni phase). The SHS with high heating rate (approx. 200 and 300 K min−1) was applied. Because the SHS product is very porous, it was milled in vibratory disk milling and consolidated by SPS technique at temperatures of 1173 K, 1273 K, and 1373 K (900 °C, 1000 °C, and 1100 °C). The microstructures of samples prepared by SHS reaction and combination of SHS reaction, milling, and SPS consolidation are compared. The changes in microstructure with increasing temperature of SPS consolidation are observed. Mechanical properties are tested by hardness measurement. The way to reduce the amount of Ti2Ni phase in structure is leaching of powder in 35 pct hydrochloric acid before SPS consolidation.
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Our previous works on the synthesis of intermetallics by Self-propagating High-temperature Synthesis revealed strong dependence of microstructure of the products on heating rate. In this paper, the application of various heating regimes and sources was tested in preparation of NiTi shape memory alloy. It was found that high heating rate (approx. over 100 °C min−1) was required to obtain the material with maximized amount of NiTi shape memory phase and no unreacted metals. Heating in electric resistance furnace preheated to the process temperature or induction heating furnace seemed to be promising for this purpose. On the other hand, Spark Plasma Sintering was found to be inapplicable, because the strongest increase of the temperature occurred on the surface of the particles, producing layers of intermetallics that further acted as diffusion barriers.
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This study presents the effects of solution annealing and subsequent aging on shape memory response of Ni-rich Ni50.8Ti49.2 alloys fabricated by Selective Laser Melting. After solutionizing, samples were heat treated at selected times at 350 °C and 450 °C and their shape memory effect, superelasticity, and transformation temperatures were determined. It was found that transformation temperatures, transformation behavior, strength, and recoverable strain are highly heat treatment dependent. Samples aged at 350 °C showed better recovery in superelasticity tests where 350 °C-18 h aged samples exhibited almost perfect superelasticity with 95% recovery ratio with 5.5% strain in the first cycle and stabilized superelasticity with a recoverable strain of 4.2% after 10th cycle. Meanwhile, 450 °C-10 h aged sample exhibited 68% recovery with a recoverable strain of 4.2% in the first cycle and stabilized recoverable strain of 3.8% after 10th cycle.