P. NOVÁK et al.: FORMATION OF Ni-Ti INTERMETALLICS DURING REACTIVE SINTERING AT 800–900 °C
679–685
FORMATION OF Ni-Ti INTERMETALLICS DURING REACTIVE
SINTERING AT 800–900 °C
OBLIKOVANJE NiTi INTERMETALNIH ZLITIN MED
REAKTIVNIM SINTRANJEM PRI 800–900 °C
Pavel Novák1, Vladimír Vojtìch1, Zuzana Pecenová1, Filip Prù{a1, Petr Pokorný1,
Davy Deduytsche2, Christophe Detavernier2, Adriana Bernatiková1,
Pavel Salvetr1, Anna Knaislová1, Kateøina Nová1, Lucyna Jaworska3
1University of Chemistry and Technology, Department of Metals and Corrosion Engineering, Technická 5,
166 28 Prague 6, Czech Republic
2Ghent University, Department of Solid State Sciences, Krijgslaan 281, S1 9000 Gent, Belgium
3Institute of Advanced Manufacturing Technology, 37a Wroclawska St., 30-011 Krakow, Poland
panovak@vscht.cz
Prejem rokopisa – received: 2016-08-17; sprejem za objavo – accepted for publication: 2016-11-16
doi:10.17222/mit.2016.257
In this work the formation of intermetallics in the Ni-Ti system by reactive sintering at 800–900 °C was studied. The mechanism
and kinetics of the reactions, which led to Ni-Ti phases, were determined by thermal analysis, in-situ XRD and the application
of an experimental model consisting of nickel-plated titanium. It was found that the formation of Ni-Ti phases below the
transformation temperature of titanium is controlled by diffusion. Above this temperature, the reactions switch to the rapid
Self-propagating High-temperature Synthesis (SHS) mode.
Keywords: reactive sintering, powder metallurgy, NiTi
V delu je bil raziskan nastanek intermetalnih zlitin v sistemu NiTi pri reaktivnem sintranju na 800-900 ° C. S termi~no analizo,
XRD-in situ analizo in uporabo eksperimentalnega modela, nikljanega s titanom, sta bila dolo~ena mehanizem in kinetika
reakcij, ki sta vodila k NiTi fazam. Ugotovljeno je bilo, da je tvorba NiTI faze pod transformacijsko temperaturo titana,
nadzorovana z difuzijo. Nad to temperaturo se reakcije spremenijo na hitro rasto~i temperaturno -sintezni na~in (SHS).
Klju~ne besede: reaktivno sintranje, metalurgija prahov, NiTi
1 INTRODUCTION
The Ni-Ti alloy called nitinol, in approximately equi-
molar proportions, is the most widely known shape-
memory alloy. The shape-memory effect in this alloy is
connected with the transformation between high-tempe-
rature cubic austenite and low-temperature monoclinic
martensite.1,2 For the practical application of these alloys,
superelasticity is very important. This phenomenon
occurs when the NiTi alloy is deformed slightly above
the martensite  austenite transformation temperature.
Deformation induces the formation of the martensite
phase, which is continuously transformed to austenite
during unloading. Due to this phenomenon, this alloy
behaves like an enormously elastic material.1,2 In
addition, the NiTi alloy is also a corrosion-resistant
material.3 Due to its exceptional properties, the NiTi
alloy is applied in both medical (dental implants, stents,
scaffolds)4,5 and technical applications (actuators, robo-
tics, etc.).6–8
The most commonly applied techniques in the indus-
trial production of nitinol alloy are melting metallurgy
processes – vacuum induction melting (VIM) and
vacuum arc remelting (VAR).9,10 In the VIM of Ti-con-
taining alloys there is a serious danger of a strong
contamination of the melt due to the high reactivity of
molten titanium.11 The VAR technique makes it possible
to prepare alloys of higher purity, but there is a problem
with homogeneity. To obtain a sufficiently homogenous
product, the VAR process has to be repeated even more
than 4 times.10 This implies that it is costly and relatively
problematic to obtain a NiTi shape-memory alloy. If a
simple production technology would be developed, the
NiTi alloy could be more frequently applied, not only in
specific areas requiring the shape-memory effect, but
also in other technical branches as a corrosion-resistant
alloy. It will be beneficial for European economy,
because this alloy does not contain any elements listed as
critical raw materials.12
A promising alternative to melting metallurgy
production routes is powder metallurgy (PM). A simple
non-conventional PM production technology is reactive
sintering. In general, the reactive sintering is a densifi-
cation process, where initial components in powder form
are transformed to a compact product via thermally-acti-
vated chemical reactions.13 These reactions are mostly
exothermic when intermetallics are formed. The route
from powders to the compact usually contains powder
blending, cold pressing and sintering.13,14 When pure
powders and an efficient protective atmosphere are
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Materiali in tehnologije / Materials and technology 51 (2017) 4, 679–685 679
UDK 621.763:621.762.3:620.172.22 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 51(4)679(2017)
applied, a high-purity product is obtained due to the
limited contact of the compressed reaction mixture with
the crucible or support during the process.
The reactive sintering process of many intermetallics,
including the Ni-Ti alloy, proceeds in two stages: a
lower-temperature diffusional stage and a rapid high-
temperature process called Self-propagating High-tem-
perature Synthesis (SHS).13–15
In the case of the reactive sintering of NiTi alloy, the
following mechanism is proposed in the literature.
During heating to 900 °C, the slow diffusional formation
of three intermetallics (Ti2Ni, NiTi and Ni3Ti) proceeds.
In our previous paper we found that traces of the Ti2Ni
phase start to form already at 471 °C, being followed by
NiTi and Ni3Ti at 632 °C.16 The transformation of -Ti to
-Ti takes place as the temperature increases and after
that the -Ti rapidly saturates with nickel. When the
temperature exceeds 942 °C, a liquid phase is formed
from the solid solution of nickel in -Ti by eutectic
transformation and thus ignites the SHS reaction.15
However, in our previous paper about the optimization of
the conditions for the reaction synthesis of the NiTi
shape memory alloy the structure corresponding to the
SHS reaction was observed already after heating to
900 °C, being very similar to the material processed at
1100 °C.17
This paper aims to explain how it is possible to
initiate the SHS reaction in NiTi alloy at 900 °C. To
prove it, the mechanism and kinetics of the Ni+Ti reac-
tions were studied by thermal analysis, in-situ XRD
analysis and an experimental model.
2 MATERIALS AND METHODS
As mentioned above, the structure appearing to result
from SHS reaction was observed in the NiTi alloy
already at 900 °C using a heating rate of approximately
300 °C/min,17 even though other references state that the
SHS reaction is triggered by the melt formation by
eutectic reaction at 942 °C. To prove the possibility to
initiate the SHS reaction under these conditions and to
describe the reaction mechanism, the following steps
were carried out:
• Thermal analysis during heating from room tem-
perature to approximately 1200 °C with the heating
rate of approximately 300 °C/min. The heating rate
was the same as in our previous paper, where the
structure corresponding to the SHS reaction was
observed after the process at 900 °C.17
• In-situ XRD analysis during heating from the labor-
atory temperature to 900°C with the heating rate of
60 °C/min. The heating rate was limited by the capa-
bilities of the device.
• Description of mechanism and kinetics of the Ni-Ti
phases’ formation at 800 °C and 900 °C using an
experimental model.
• Reactive sintering of Ni-Ti samples at 800 °C and
900 °C with the heating rate of 300 °C/min.
Experimental material for thermal analysis and
in-situ XRD was prepared by blending of nickel powder
(particle size < 150 μm, > 99.99 % purity, supplied by
Aldrich) and titanium powder (particle size <44 μm, >
99.5 % purity, supplied by Alfa Aesar). Green bodies
with a cylindrical shape of 10 mm in diameter and
approximately 5 mm in height were prepared by uniaxial
cold pressing of the powder blends under a pressure of
630 MPa using LabTest 5.250SP1-VM universal loading
machine.
Thermal analysis was carried out during heating in
the induction furnace from room temperature to 900 °C
with the heating rate of approximately 300 °C/min. An
optical pyrometer (Optris P20 2M) was used to record
the temperature profile of the reaction.
In-situ XRD analysis (Cu-K radiation) was carried
out during heating from laboratory temperature to 900 °C
in a helium atmosphere with a heating rate of 60 °C/min.
An experimental model was established and success-
fully applied in our previous works in the Fe-Al,
Fe-Al-Si and Ti-Al-Si systems.18,19 In the mentioned
papers, the experimental model consisted of an iron or
titanium sample submerged in molten aluminium or
Al-Si alloy. The aim of these experiments was to des-
cribe the kinetics of the formation of intermetallics in
these particular alloys systems, because directly in
reactive sintering process it is not possible. In the present
work a model (Ni-Ti diffusion couple) consisting of a
titanium bulk sample coated with nickel was applied.
The nickel-plated titanium simulates the interaction
between the compressed titanium and nickel powder
particles. The galvanic nickel plating was carried out at
60 °C in a Watts’s bath containing 300 g/L NiSO4.7H2O,
40 g/L NiCl2 and 40 g/L H3BO3 using a current density
of 10 A/m to the final thickness of approximately 20 μm.
Before galvanic plating, titanium samples were sand-
blasted with alumina for 10 min, pickled in 35 % HCl at
60 °C for 12 min and activated in 35 % HCl at 20 °C for
15 s. The coated samples were annealed for 30–180 min
in evacuated and sealed silica ampoules at 800 °C. At
900 °C, the experiment was stopped after 120 min,
because the whole thickness of the nickel coating reacted
to form Ni-Ti phases. The heating rate during the experi-
ments was approximately 300 °C/min. The phase
composition of the obtained multilayer systems was
identified by the X-ray diffraction (XRD) method using
PANalytical X’Pert Pro diffractometer (Cu-K radiation).
PANalytical X’Pert HighScore Plus software with the
PDF-2 database was used to process and to evaluate the
XRD patterns. The microstructure of intermetallics’
layers was examined with a TESCAN VEGA 3 scanning
electron microscope equipped with OXFORD Instru-
ments X-max EDS SDD 20 mm2 detector (SEM-EDS).
Samples were mechanically ground, polished and etched
using modified Kroll’s reagent (10 mL HF, 40 mL HNO3
P. NOVÁK et al.: FORMATION OF Ni-Ti INTERMETALLICS DURING REACTIVE SINTERING AT 800–900 °C
680 Materiali in tehnologije / Materials and technology 51 (2017) 4, 679–685
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
and 50 H2O) before the microstructure observation.
Image analysis was carried out by the means of ImageJ
software in order to determine the thickness of the
intermetallics’ layers. A process controlling the forma-
tion of intermetallics was determined by fitting the layer
thickness with a parabolic growth equation. When a
process is controlled by diffusion of species through a
reaction product, it is generally described by the para-
bolic law, written as Equation (1):
d k t= ⋅p (1)
where d is the layer thickness (μm) and t represents the
annealing duration (s).20
3 RESULTS AND DISCUSSION
The temperature profile during heating of the
compressed powder mixture of nickel and titanium was
recorded using an optical pyrometer. During heating in
an induction furnace by the rate of nearly 300 °C/min,
the temperature increased almost linearly up to
approximately 880 °C (Figure 1). After that, the increase
of the slope, i.e., the heating rate, can be observed. This
increase can be probably be attributed to the initiation of
the strongly exothermal SHS reaction. This temperature
corresponds well with the transformation of titanium
from hexagonal (hcp) structure to cubic (bcc) one.21
After achieving approximately 940 °C, the slope of the
curve increases even more rapidly, probably due to the
formation of a liquid phase by the eutectic transforma-
tion in the Ni-Ti system (942 °C).21 The liquid medium
supports the SHS reaction by increasing the diffusion
rate of the reactants. The maximum temperature
achieved by the reactions was 1137 °C. This indicates
that the sample was partially molten due to eutectic
reactions at 942 °C and 1118 °C, but did not exceed the
melting point of the NiTi phase (1310 °C).21 It can be
expected that the temperature at the reaction front can be
even higher, but the heat is quickly transferred to the rest
of the sample and therefore the overall temperature of
the sample is lower.
The in-situ XRD analysis during heating of the
compressed powder mixture from the room temperature
to 900 °C with a rate of 60 °C/min (Figure 2) shows the
continuous shift of the diffraction lines of nickel and
titanium to lower angles with increasing temperature.
This phenomenon is caused by the thermal expansion of
the powders. In addition, the phase transformation of
hexagonal -Ti to cubic -Ti can be observed around the
temperature expected according to the Ni-Ti phase
diagram (882 °C).21 In addition to this phase transforma-
tion, the weak diffraction lines of the Ti2Ni (cubic
structure Fd3m)21 and NiTi (austenite, cubic structure
Pm3m)21 phases start to be visible at approximately
750 °C. The intensity of the lines of the Ti2Ni phase
increase very rapidly, when a temperature of approxi-
mately 880 °C is achieved (Figure 2). This increase is
immediately followed by the rapid formation of NiTi
(austenite, cubic structure Pm3m)21 and Ni3Ti (hexagonal
P63/mmc)21. It implies that the SHS reaction between
nickel and titanium is initiated by the transformation of
titanium to its cubic allotropic modification. It forms the
Ti2Ni phase, which then reacts with nickel and/or the
Ni3Ti phase to form NiTi (austenite). The same phases’
formation sequence was also observed for the diffusion
stage of the reactive sintering process during long-term
annealing at 500–650 °C. It implies that Ti2Ni forms
preferentially and it cannot be avoided in both
low-temperature diffusion process, as well as during
SHS mode.
To be able to describe the kinetics of the formation of
intermetallics, the experimental model simulating the
interaction between compressed nickel and titanium
particles was applied. The model consisted of a bulk
titanium sample covered by a nickel layer of approxi-
mately 20 μm in thickness. The model samples were
P. NOVÁK et al.: FORMATION OF Ni-Ti INTERMETALLICS DURING REACTIVE SINTERING AT 800–900 °C
Materiali in tehnologije / Materials and technology 51 (2017) 4, 679–685 681
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 2: Variation of XRD intensities of NiTi46 (in mass fractions,
(w/%)) compressed powder mixture during heating from room
temperature to 900 °C (heating rate of 60 °C min–1)
Figure 1: Heating curve of Ni-Ti compressed powder mixture
(heating rate of approx. 300 °C min–1, recorded by optical pyrometer)
annealed at 800 °C and 900 °C and the thickness of
formed layers of intermetallics was measured to describe
the difference between in the mechanism and kinetics of
the Ni+Ti reactions below and above the    trans-
formation temperature in titanium. During annealing at
800 °C, the layers of Ti2Ni, NiTi and Ni3Ti are formed
between the titanium and nickel coating (Figure 3a to
3d). The presence of these phases was confirmed by the
XRD (Figure 4) and local EDS analysis (Table 1).
While the XRD analysis revealed all present layers on
the sample, the EDS analysis was used to identify the
sequence of the layers between the reacting titanium and
nickel. The dependences of the thickness of the Ti2Ni
and NiTi layers on the process duration exhibit the
parabolic shape (Figure 5). The calculated parabolic rate
constants for all the samples indicate that the growth of
these layers slightly slows down with the process
duration (Table 2). The layer of Ni3Ti (hexagonal
P63/mmc)21 phase starts to grow after 60 min of
annealing and then it follows the parabolic law. The
formation of the Ni3Ti phase is probably the reason why
the growth of the NiTi and Ti2Ni layers decelerates
during longer annealing. As the Ni3Ti phase grows, NiTi
and/or Ti2Ni are consumed, as well as nickel. The results
confirm that during the diffusion stage, the process starts
with the formation of Ti2Ni and NiTi layers, followed by
Ni3Ti after longer annealing duration. Almost the same
shape of the kinetic curves was also observed during
annealing at 650 °C in our previous paper, where the
formation of Ni-Ti intermetallics was studied during
long-term annealing at 500-650 °C.16
Table 1: Chemical composition (EDS) of layers observed after
annealing of model samples at 800 °C for 120 min
Phase
Content (in amount fractions, (a/%))
Ti Ni
Ni 4.0±0.2 96.0±0.2
Ti2Ni 65.7±0.6 34.3±0.6
NiTi 51.4±0.4 48.6±0.4
Ni3Ti 27.9±1.8 72.1±1.8
Ti 98.4±0.2 1.6±0.2
Annealing at 900 °C leads to the layers of Ti2Ni and
NiTi (Table 3, Figure 4), which grow slightly more
rapidly up to 60 min (Figure 6a) than at 800 °C (Fig-
ure 5). However, the thickness of the NiTi phase layer is
very non-uniform. After 60 min, the thick layer of Ni3Ti
(hexagonal P63/mmc) phase arises in the structure. The
P. NOVÁK et al.: FORMATION OF Ni-Ti INTERMETALLICS DURING REACTIVE SINTERING AT 800–900 °C
682 Materiali in tehnologije / Materials and technology 51 (2017) 4, 679–685
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 3: Microstructure of Ni-Ti model samples annealed at 800 °C
for: a) 30 min, b) 60 min, c) 120 min and d) 180 min
Figure 5: NiTi, Ti2Ni and Ni3Ti layer thickness on model samples vs.
duration of annealing at 800 °CFigure 4: XRD patterns of model samples annealed at 800 °C and 900 °C
Ni3Ti layer consumes almost the whole residual thick-
ness of the nickel coating (Figure 6b). The Ni3Ti layer is
very porous, which confirms the previous results, where
the Ni3Ti phase was found to be a source of porosity.16
When continuing the annealing, the Ni3Ti phase
disappears completely, since it reacts with Ti2Ni in order
to form a thick layer of NiTi phase (Figure 6c). The
experiment was stopped after 120 min, because the
whole layer was composed of NiTi and Ti2Ni phases
(Figures 6c and 7). Calculated parabolic rate constants
(Table 4) show that the formation of NiTi (austenite) and
Ni3Ti is not diffusion-controlled at 900 °C, because the
constants vary strongly with the process duration. It con-
firms that the process runs in SHS mode.
Table 4: Parabolic rate constant of the formation of Ti2Ni, NiTi and
Ni3Ti layer vs. annealing duration at 900 °C
Annealing
duration
(min)
Parabolic rate constant of the growth of
layers (×10–4 μm s–1)
Ti2Ni NiTi Ni3Ti
30 16.1 46.7 0
60 87.1 40.1 267.0
120 50.0 325.0 0
To prove the results of the experimental model, the
real compressed powder mixtures were prepared and
heated at 800 °C and 900 °C with a holding time of 120
min. The same heating rate as in the experimental model
(approx. 300 °C/min) was applied. The sample prepared
by heating at 800 °C is composed of unreacted nickel
and titanium particles covered by layers of Ti2Ni, NiTi
and Ti3Ti of very similar thickness like in the model sys-
tem (Figure 8a). On the other hand, the sample react-
ively sintered at 900 °C is comprised of Ti2Ni particles in
a NiTi matrix. The morphology of the sample changed
from cylindrical shape to irregular shape due to partial
melting during the process (Figure 9). This indicates that
P. NOVÁK et al.: FORMATION OF Ni-Ti INTERMETALLICS DURING REACTIVE SINTERING AT 800–900 °C
Materiali in tehnologije / Materials and technology 51 (2017) 4, 679–685 683
MATERIALI IN TEHNOLOGIJE/MATERIALS AND TECHNOLOGY (1967–2017) – 50 LET/50 YEARS
Figure 7: NiTi, Ti2Ni and Ni3Ti layer thickness on model samples vs.
duration of annealing at 900 °C
Figure 6: Microstructure of Ni-Ti model samples annealed at 900 °C
for: a) 30 min, b) 60 min and c) 120 min
Table 2: Parabolic rate constant of the formation of Ti2Ni, NiTi and
Ni3Ti layer vs. annealing duration at 800 °C
Annealing
duration (min)
Parabolic rate constant of the growth of
layers (×10–4 μm s–1)
Ti2Ni NiTi Ni3Ti
30 8.0 43.6 0
60 6.3 36.0 11.1
120 4.5 24.5 17.0
180 3.3 19.6 17.1
Table 3: Chemical composition (EDS) of layers observed after
annealing of model samples at 900 °C for 120 min
Phase
Content (in amount fractions, (a/%))
Ti Ni
Ti2Ni 66.8±0.5 33.2±0.5
NiTi 53.2±0.9 46.8±0.9
Ti 97.8±0.3 2.2±0.3
during annealing of the powder mixture at 900 °C,
enormous heat is evolved due to the SHS reaction.
The above-presented results show that the SHS
reaction can be initiated immediately after exceeding the
   transformation temperature of titanium. It means
that it is not necessary to achieve the melt formation by a
eutectic reaction at 942 °C. In addition, the rapid SHS
process is supported significantly when a high heating
rate is applied.17 During rapid heating, the diffusion is
strongly suppressed. Therefore, the SHS can be initiated
by the reaction of solid particles of nickel and titanium,
having a cubic structure, producing a mixture of cubic
Ti2Ni and NiTi phases at the end of the process.
4 CONCLUSIONS
In this work the mechanism and kinetics for the for-
mation of Ni-Ti phases at 800 °C and 900 °C were
studied. The results revealed that the rapid SHS reaction
is initiated in the solid state by the → phase transfor-
mation of titanium. Below this temperature, the growth
of Ni-Ti intermetallics is controlled by diffusion of the
reactants. Above this temperature, the rapid SHS
reaction is initiated. It was proved with an experimental
model, where a significant change in the rate constant of
the formation of Ni-Ti intermetallics was observed at
900 °C, compared with 800 °C. It was confirmed by the
change of the slope of the heating curve in thermal
analysis, as well as by the in-situ XRD analysis, where a
rapid increase of the intensity of the diffraction lines of
Ni-Ti intermetallics was observed after exceeding
approximately 880 °C. The SHS process starts with the
formation of Ni3Ti phase, which reacts with Ti2Ni
already present from diffusion stage and forms the NiTi
shape-memory phase. The temporary formation of Ni3Ti
phase results in an increase of the materials’ porosity.
Acknowledgement
This research was financially supported by Czech
Science Foundation, project No. 14-03044S. Partial
support by COST Action CA15102 is also greatly appre-
ciated.
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