EFFECT OF HEAT TREATMENT ON THE SUPERELASTICITY AND
HARDNESS OF NITI
Omar Mortagy and Mahmoud Farag
omortagy@aucegypt.edu and mmfarag@aucegypt.edu
Science and Technology Research Center and
Mechanical Engineering Department
The American University in Cairo
P.O. Box 2511, Cairo 11511, Egypt
Fax: (20-2)792 3352
1. ABSTRACT
This paper investigates the effect of heat treatment on grain size, microstructure,
superelasticity and mechanical properties of severely deformed NiTi alloy (51 atomic% Ni).
The as received alloy, which was a 25 mm bar fully annealed at 850˚C with austenitic
structure, was first subjected to severe plastic deformation by machining on a lathe under
water cooling with cutting speeds of 0.118, 0.59 and 1.18 m/s, depth of cut of 2 mm and feed
of 0.09 mm/rev. The machined chips were then annealed for one hour at 200, 300, 400, 500,
600, 700, and 800˚C in order to obtain different microstructures. SEM, EDX and optical
microscopy were used to study the microstructure, an MTS nanohardness machine with a
Berkovich indenter was used to measure superelasticity and nanohardness, and a Vickers
microhardness tester with a load of 4.9 N was used to measure the microhardness of the
material. The results of this work showed that the behavior of the NiTi alloy is influenced by
cold work and heat treatment. A qualitative model based on the density of dislocations and
precipitates was used to explain the results. Relatively high superelasticity and microhardness
were exhibited by the material after cold work followed by low temperature annealing for one
hour. In spite of the fact that microhardness readings were influenced by elastic recovery, the
results showed that it was more sensitive to changes in the microstructure than nanohardness.
Keywords: NiTi, machining, annealing, precipitates, superelasticity, microhardness,
nanohardness
-1-
2. INTRODUCTION
Improving superelastic and shape memory properties of Ni-rich NiTi alloys is a very
challenging and important subject due to the commercial importance of this alloy. Aging is a
very effective way of improving shape memory and superelastic properties [1-4]. Aging
causes precipitates, such as Ti11Ni14, Ti3Ni4, Ti2Ni3, and TiNi3 to form, which strengthen the
matrix and hence improve superelasticity [5]. Ti11Ni14 precipitates were shown to form when
aging a Ti-51 at.% Ni alloy at 300, 400, 500, and 600ºC for one hour [1, 2]. Upon aging at
300 and 400ºC, fine Ti11Ni14 particles were formed in the matrix and the material exhibited
shape memory behavior. At 500ºC Ti11Ni14 precipitates had grown in size but were still
coherent and the material also showed shape memory behavior. No precipitates were observed
at 600ºC, and, as a result the alloy did not exhibit the shape memory property [2]. Ti3Ni4
precipitates were also reported to form upon aging a Ti-50.9 at.% Ni alloy [3]. However, the
precipitates did not form at low temperatures, e.g. 200ºC , except after 100 hours of aging,
whereas at 400ºC they formed after aging for 5 minutes [3]. The size, density, and coherency
of the Ti3Ni4 precipitates were important regarding improvement of superelastic properties. It
was shown that the increase in the density of fine Ti3Ni4 precipitates was more effective in
hindering dislocation motion than the coarsening of these precipitates and thus caused the
superior shape memory properties [3, 4]. Coherency of precipitates had significant influence
on shape memory behavior [2] and on shape memory properties, such as martinsitic
transformation stress [6].
Another strategy to improve superelasticity is to widen the superelastic temperature
range and center it around the application temperature. Increasing the austenitic strength, or in
other words the critical stress for slip, by cold working is a very effective way for increasing
the width of the superelastic temperature window. Cold working, however, must be followed
by an annealing process to preserve the ability to stress induce martensite. It was reported that
cold-rolling induced dislocations strengthened the matrix of NiTi and acted as obstacles to
slip [7]. Twin boundary movement, which is the fundamental deformation mechanism of
stress induced martensite (SIM), was suggested to be less affected by the presence of the
dislocations, and thus twin boundaries were able to move easily to produce the SIM. Because
the irreversible slip was depressed by the dislocations during the SIM transformation,
pseudoelasticity was improved [7]. Low temperature annealing for one hour after cold work
was more effective at improving superelastic response than high temperature annealing [8].
At low temperatures dislocations were thermally arranged and a high density of fine
precipitates were present. This combination resulted in a significant increase in the critical
stress for slip and thus improved superelasticity [8].
Hardness measurements of superelastic NiTi received some attention in recent years
[6, 9]. It was recommended to use nano-indentation instead of micro-indentation with
superelastic NiTi in order to avoid the problem of elastic recovery and yield true hardness
values [9]. It was also noted that hardness re-increased after higher temperature annealing due
to precipitation dissolution [6].
Superelastic measurements of NiTi can be made using nano-indentation [10, 11].
Although it is recommended to use spherical tips when measuring superelasticity to avoid
plastic deformation [10, 11], some conclusions on superelastic behavior can be drawn if a
Berkovich indenter is used [11].
-2-
Many different techniques have been used to produce fine-grained NiTi. These
techniques include mechanical alloying [12], rapid solidification [13], plastic deformation
[14], powder methods [15], and machining [16-20]. Machining, which is the method used in
this research has many advantages: it is a cheap process, includes very large strains and strain
rates, is a one-stage deformation process, can handle high-strength metals and alloys, and is
supposed to produce ultra-fine structured chips [16].
The purpose of this research is thus to study the effect of annealing on microstructural,
mechanical and superelastic properties of NiTi chips produced by machining.
3. EXPERIMENTAL TECHNIQUES
The material used in this work is austenitic Nickel Titanium 56 wt% Ni-Ti (51 at%
Ni). The material was supplied by Euroflex GmbH as a 25 mm diameter rod fully annealed at
850ºC.
The material was subjected to severe plastic deformation by machining on a center
lathe. The feed rate was 0.09 mm/rev, depth of cut 2 mm, and the cutting speed was varied
between 0.118 and 1.18 m/s. Cooling was used during cutting. The resulting chips were
annealed for 1 hour at 200, 300, 400, 500, 600, 700, and 800ºC. After annealing the chips
were quenched in water at room temperature.
Microstructural investigations were carried out using optical microscopy and a LEO
55 field emission scanning electron microscope (FESEM). Following grinding and polishing,
chips were etched using an etching solution of composition HF + HNO3 + H2O (1:4:5). Grain
size measurements were made using the line intercept method. Energy dispersive x-ray
(EDX) analysis was used to determine the compositions of any existing phases in particles
and grain interiors. The EDX was part of the FESEM. At each selected point usually 5 EDX
readings were taken.
An MTS Nanoindenter XP with a Berkovich indenter was used to measure
superelasticity and nanohardness. Superelastic measurements were carried out after polishing
the surface using alumina powder with particle size 1 μm followed by 0.05 μm. This was
necessary to avoid any errors due to surface roughness. A Berkovich indenter with a load of
20 mN was used in this process. The maximum and residual depths reached by the indenter
for each test were recorded. Using these values the recovered indentation depth was
calculated for each test by subtracting the residual depth from the maximum depth. Finally, an
average value of all recovered indentation depths was determined for the sample.
Nanohardness measurements were performed also using a Berkovich indenter. The
depth control option was used to fix indentation depth at 1000 nm. The applied loads ranged
from 45 to 75 mN. A range of 25 to 40 points per sample were tested and the average was
taken. Hardness values were given in GPa.
A Vickers microhardness tester was used to measure microhardness of the samples. A
load of 4.9 N and testing duration of 15 seconds were selected as the loading parameters. The
loading step took 5 seconds and then the load was held at the specific value for 10 seconds
before unloading started. An average of 21 points per sample was tested.
-3-
4. RESULTS & DISCUSSION
Microstructure
Figure 1 shows that the as-received material consisted of equi-axed grains of
approximate size 25.5 µm, measured by the intercept method. Some grain interiors have
structures, which seem to be of different phases. Some of these structures are arranged in a
twin arrangement.
Fig. 1: SEM image showing microstructure of as-received NiTi
a) as-received
b) Vc = 0.118 m/s
c) Vc = 0.59 m/s
d) Vc = 1.118 m/s
Fig. 2: Optical microscope images showing microstructures of as-received NiTi and as-machined chips at
3 cutting speeds
-4-
Effect of machining on microstructure
After machining under cooling the original material was severely deformed. This is
shown in figures 2(b) and (c). Generally, the grains were elongated in the direction of
machining, but in most areas a distinct separation between the grains could not be seen due to
the severe deformation. Grain interiors were heavily deformed as shown by the arrows.
In addition, no apparent difference was detected between the microstructures resulting
from cutting at speeds 0.118 m/s and 0.59 m/s. At the higher speed of 1.18 m/s, however, the
resulting microstructure was completely different. As illustrated in figure 2(d), there were
very clear shear bands, and the grains between shear bands were almost undeformed. Shear
bands are narrow bands running at the bottom of the chip and branching off to the chip teeth.
They are characterized by a high degree of instability and intense shear strains, which are
much higher than the shear strains in the rest of the material.
Effect of annealing on microstructure
Figure 3 shows images taken from heat treated samples.
a) T = 200ºC
b) T = 300ºC
c) T = 400ºC
d) T = 500ºC
Fig. 3: Chip microstructures cut at 0.118 m/s and annealed (a-e SEM; f & g optical microscope)
-5-
e) T = 600ºC
f) T = 700ºC
g) T = 800ºC
Fig. 3: (continued)
No obvious change in microstructure appeared when the chips were annealed between
200ºC and 500ºC (fig. 3a-d). Grains remained elongated and internally deformed. At 600ºC
(fig. 3e), however, small equi-axed grains had formed indicating that recrystallization had
taken place at a temperature between 500ºC and 600ºC. Grain growth followed at 700ºC (fig.
3f) and continued at 800ºC (fig.3g). The size of the recrystallized grains versus heat treatment
temperature is plotted in figure 4.
Grain size [microns]
30
25
y = 0.0933x - 51.83
20
Heat-treated
15
as-received
10
Linear trend line
5
0
500
600
700
800
900
Temperature [C]
Fig. 4: Variation of grain size with annealing temperature
The variation of grain size with heating temperature is almost linear. A best fit equation can
be written to relate grain size with temperature:
-6-
G.S. = 0.0933T – 51.83
where G.S. = grain size in [μm]
and
T = temperature in [ºC] between 600 and 800ºC
Phase distribution
Figures 5(a) & (b) are plots of the atomic percentage of Ni found in particles and grain
interiors, respectively, versus annealing temperature. As-received NiTi was included for
comparison. The vertical lines represented the compositions of the various compositions,
which were found in literature and which were considered to have an effect on shape memory
and superelastic properties. In addition, a vertical line representing the average as-received
composition, i.e. 51% Ni, was added for easy reference.
Ti11Ni14
Ti2Ni
Ti2Ni3
TiNi3
900
TiNi
Ti3Ni4
(a)
800
as-received
T= 200 C
T = 300 C
T = 400 C
T = 500 C
T = 600 C
T = 700 C
TiNi
Ti2Ni
Ti11Ni14
Ti3Ni4
Ti2Ni3
TiNi3
700
Temperature [C]
600
500
400
300
200
100
0
25
30
35
40
45
50
55
60
65
70
75
80
Ni - atomic %
Ti2Ni
TiNi
Ti11Ni14 Ti2Ni3
900
TiNi3
Ti3Ni4
(b)
800
T = 200 C
T = 300 C
T = 400 C
T = 500 C
T = 600 C
T = 700C
T = 800C
as-received
Ti2Ni
Ti11Ni14
Ti3Ni4
Ti2Ni3
TiNi3
TiNi
700
Temperature [C]
600
500
400
300
200
100
0
30
35
40
45
50
55
60
65
70
75
80
Ni - atomic %
Fig. 5: Atomic % of Ni in particles (a), and grain interiors (b) at the different temperatures
Figure 5(a) shows that the larger particles in the matrix were mainly of the Ti2Ni
composition. According to literature findings, these particles did not have any role in
-7-
influencing mechanical or superelastic properties. They were, therefore, not investigated
further.
To simplify the analysis and understanding of figure 5(b), points were grouped into
two Ni-atomic % groups, namely matrix and precipitates. The matrix contained all
compositions at and below 53% Ni and precipitates contained all compositions starting from
54% Ni and above. The most important precipitates that were expected were Ti11Ni14 and
Ti3Ni4, which had Ni compositions of 56 and 57%, respectively. It can be seen that these two
compositions were very close and difficult to differentiate between them. If one considered a
2% fluctuation, then the precipitates group would range from 54 to 59%. The matrix was
considered to be any composition below 54% Ni with the majority of the points lying between
48 and 53% Ni. Frequency distributions of the heat treated samples were then plotted to
determine the percentage of occurrence of each phase at each condition. The plots are shown
in figure 6.
Frequency distribution - T = 200 C
100
90
80
70
%
60
50
200
40
30
20
10
0
Matrix
Precipitates
Ni - atomic %
Frequency distribution - T = 300 C
Frequency distribution - T = 400 C
100
90
90
80
80
70
70
60
60
300
%
%
50
50
400
40
40
30
30
20
20
10
10
0
0
Matrix
Precipitates
Matrix
Ni - atomic %
Precipitates
Ni - atomic %
Frequency distribution - T = 500 C
Frequency distribution - T = 600 C
100
90
90
80
80
70
70
60
60
500
%
%
50
50
600
40
40
30
30
20
20
10
10
0
0
Matrix
Precipitates
Matrix
Ni - atomic %
Precipitates
Ni - atomic %
Fig. 6: Frequency plots of Ni composition groupings for each condition
-8-
Frequency distribution - T = 700 C
Frequency distribution - T = 800 C
100
100
90
90
80
80
70
70
60
50
700
%
%
60
50
40
40
30
30
20
20
10
800
10
0
0
Matrix
Precipitates
Matrix
Precipitates
Ni - atomic %
Ni - atomic %
Fig. 6: continued
The above frequency plots can be summarized in figure 7, which showed the
compositional variation of each of the 2 phase groups with annealing temperature. Trend lines
have been added.
100.0
95.0
90.0
85.0
80.0
75.0
Frequency of occurrence %
70.0
65.0
60.0
Matrix
Precipitates
Linear (Matrix)
Poly. (Precipitates)
55.0
50.0
45.0
40.0
35.0
30.0
25.0
20.0
15.0
10.0
5.0
0.0
200
300
400
500
600
700
800
Temperature [C]
Fig. 7: Variation of grain interior phase composition with annealing temperature
From figures 6 and 7 it was obvious that the dominating phase in the grain interiors for
all conditions were that of the original matrix, i.e. 51% Ni. The trend line was almost
horizontal, which indicates no significant changes in the matrix composition over the entire
temperature range.
Precipitates were, according to literature [1-3, 6], mainly of Ti11Ni14 and Ti3Ni4
compositions, and were combined together here under one group. The trend line showed an
initial increase at 200 and 400ºC then a final decrease in the range 600-800ºC. According to
literature [3], Ti11Ni14 and /or Ti3Ni4 precipitates do not form at 200ºC except after 100 hours
of aging, and formed at 300-350ºC after 1-1.5 hours of aging [1, 2, 6]. Based on these
findings, precipitation in our case did not start at 200ºC because of the relatively short
annealing time (1 hour) and had, therefore, the lowest frequency of occurrence as shown in
figure 7. As annealing temperature increased precipitation also increased as indicated by the
higher frequencies of occurrence at 300 and 400ºC. Frick et al. [6] reported that at 550ºC after
1.5 hours of annealing precipitates grew significantly and became incoherent. They also
-9-
mentioned that high dislocation densities caused Ti3Ni4 precipitates to lose coherency earlier
than non-deformed NiTi. Therefore, it can be assumed that in our case in the range of 500ºC
after 1 hour of annealing precipitates started to lose coherency. The decrease in the range 600800ºC is consistent with literature: Nishida et al. [2] reported that no precipitation was
detected after heating at 600ºC , and Frick et al. [6] reported that dissolution of precipitates
took place after heating at 600ºC for 1.5 hours. Hence, it was concluded that after annealing at
temperatures of 600ºC and above precipitates began to dissolve in the matrix causing the
decrease in volume and/or number of precipitates. This was reflected in the results of figure 7.
Microhardness
Figure 8 shows that micro-hardness values decreased as annealing temperature
increased with a large drop taking place between 500˚C and 600˚C, and then slightly
increased again at 700˚C and 800˚C. These results can be explained as follows:
 The hardness value at 200˚C showed significant increase with respect to the asreceived value. This was due to retained dislocations. The 200˚C hardness was even
higher than the as-machined value. In addition to the organization of retained
dislocations, which play an important role in increasing hardness, this observation
could be explained by considering the superelastic behavior of the heated samples. As
will be shown later, superelasticity improved after machining and subsequent lowtemperature heating at 200˚C. Microhardness is measured from the residual area after
load removal. In the case of superelastic materials the elastic part is recovered upon
load removal causing a decrease in the size of the residual area. Consequently, the
resulting hardness value will appear larger.
 At 300˚C the hardness value was still high, which indicated that there was still a
significant portion of retained dislocations. Similarly to the 200˚C sample, the 300˚C
hardness value was also larger than the as-machined one. The reason for this
improvement in hardness was due to improved superelasticity. In addition, according
to literature [2, 6], coherent precipitates are formed at 300˚C and above. Therefore, it
can be said that in our case precipitates, which formed at 300˚C and were still small
and coherent, added to the hardness of the specimen and caused this higher value.
 As the temperature increased from 400 to 500˚C, dislocations were annihilated.
Precipitates increased in size [6, 8] approaching incoherency and, hence, did not add
much to hardness any more. Hardness decreased as a result.
 There was a significant drop in hardness between 500˚C and 600˚C. This indicated
that recrystallization took place between 500˚C and 600˚C, and it is well known that
softening accompanies recrystallization due to formation of new grains free of
dislocations.
 At temperatures higher than 600˚C the material experienced a slight increase in
hardness. In fact, the hardness at 800˚C was almost the same as that of the as-received
material. Such an increase was reported in literature by Frick et al. [6]. This reincrease in hardness was explained by the solutionizing process that NiTi experienced
at high temperatures. Ti3Ni4 precipitates partially dissolved and were of similar size as
the as-received material. Hardness, therefore, approached that of as-received NiTi.
- 10 -
heat-treated chips
as-machined Vc=0.118 m/s
as-received NiTi
250
Microhardness [HV]
200
150
100
50
0
200
300
400
500
600
700
800
Heat treatment temperature [C]
Fig. 8: Variation of microhardness with annealing temperature (scatter 7%)
Nanohardness
Figure 9 shows the effect of heat treatment on the nano-hardness of NiTi chips. After
heating at 200˚C for 1 hour, average nanohardness dropped from the as-machined value and
continued its decreasing until it reached a minimum at 700˚C. At 800˚C average nanohardness
increased again approaching the as-received value. This decreasing trend of nanohardness
values with increasing annealing temperature was, similar to microhardness results, obviously
a result of dislocation rearrangement and annihilation. Loss of coherency and growth of
precipitates beyond 500˚C was also responsible to the loss of nanohardness at these
temperatures.
5
Nanohardness [GPa]
4
Hardness [GPa]
as-machined chip
3
as-received NiTi
2
1
0
0
100
200
300
400
500
600
700
800
Heat treatm ent tem perature [C]
Fig. 9: Variation of nanohardness with annealing temperature (scatter 7.6%)
- 11 -
The increase in nanohardness after annealing at 800˚C to about the as-received value,
was very similar to the increase in hardness observed in the microhardness results and can be
attributed to solutionization with precipitates becoming partially dissolved and of similar size
as the as-received material. Therefore, nanohardness increased approaching the as-received
value.
Superelasticity
Figure 10 shows the variation of the superelastic response after 1 hour of annealing at
the different temperatures. A clear improvement was seen after annealing at 200ºC. Although
the recovered depth decreased as the annealing temperature increased to 300ºC and 400ºC,
there was still an obvious improvement over the value of as-received NiTi and the asmachined NiTi chip. At 500 and 600ºC the recovered depth was close to but still slightly
higher than the value of as-received NiTi. The superelasticity of samples annealed at 700ºC
showed a sharp increase, and then decreased at 800ºC. Recrystallization, which took place
between 500 and 600˚C, had no apparent effect on superelasticity.
annealed NiTi chips
as-received NiTi
Recovered indentation depth [nm]
310
as-machined NiTi
290
270
250
230
210
190
170
150
0
100
200
300
400
500
600
700
800
900
Tem perature [C]
Fig 10: Variation of recovered indentation depth with temperature (scatter 7.8%)
These results were in agreement with previous findings [5, 7, 8], in which lowtemperature annealing following cold working improved superelastic properties. This
improvement can be attributed to 2 factors:
a) The first factor was the retained dislocations, which resulted from cold
working. As was indicated by Miyazaki et al. [4], the dislocations were
arranged by the annealing process. The arranged dislocations increased the
critical stress for slip and acted as obstacles to irreversible slip thus delaying
the onset of plastic deformation and improving superelasticity.
b) The second factor was the high density of fine precipitates formed during
intermediate temperature annealing (300-400ºC), namely Ti11Ni14 or Ti3Ni4,
which increased the critical stress for slip and hence also caused improvement
in superelasticity [8].
A proposed model that qualitatively shows the interaction between dislocations and
precipitates and their effect on superelasticity is given in figure 11.
- 12 -
Superelasticity
Effect of
dislocations
100
200
300
Effect of
precipitates
400
500
600
700
800
Temperature [ºC]
Fig. 11: Effect of dislocations and precipitates on superelasticity
The model shown in figure 11 suggests that at low annealing temperatures, e.g. 200400ºC, cold work induced dislocations exist in relatively large densities but are thermally
arranged and thus increase the critical stress for slip improving superelasticity. Above 400ºC
the effect of dislocations decreases because of their annihilation. At 600ºC recrystallization
takes place resulting in a significant decrease in dislocation density. The effect of dislocations
on superelasticity reaches a minimum and continues at this minimum at 700 and 800ºC. On
the other hand, precipitates start to form at about 300ºC where they are of small size and
coherent. As the temperature increases, in the range of 400ºC, precipitates grow in size but are
still coherent and of large density. They, therefore, have a positive effect on superelasticity.
Coarsening, incoherency and decrease in density takes place at about 500ºC [6, 8] and as a
result the effect on superelasticity decreases. Dissolution of the precipitates starts between
600 and 700ºC, and precipitates decrease in size [6]. Their effect on superelasticity reaches its
maximum at 700ºC. At 800ºC precipitates continue to dissolve in the matrix leading to a
lower effect on superelastic response. At this temperature size and distribution of precipitates
are similar to the as-received material. Therefore, recovered depth seems to be approaching
the value of as-received NiTi, which was fully annealed at 850ºC.
It is also possible that precipitates improves superelastic behavior through another
mechanism. As reported by Frick et al. [6], the transformation stress required to induce
martensite in annealed cold-drawn Ti-50.9 at.% Ni decreased as the annealing temperature
increased. The authors attributed this behavior to dislocation annihilation and Ti3Ni4
precipitates. During annealing dislocations were annihilated. In addition, limited precipitation
growth took place, which created an internal stress that had a significant influence on
promoting the matrix to change from austenite to martensite, at a lower stress. This argument
is valid at low temperatures, i.e. 300-400˚C [6].
Superelasticity and hardness
Figures 12 and 13 show the relations between recovered indentation depth and microand nanohardnesses. Generally, superelasticity was function of micro-and nanohardness
following annealing: as hardness increased, superelasticity also increased and, as shown by
the dotted trend lines in both figures. This is because superelasticity depends greatly on the
strength of the matrix, which by itself is affected by the existence of dislocations, precipitates,
or both. The 700ºC sample, however, did not seem to follow this trend: although its hardness
was low, it exhibited significant superelastic behavior. This observation could not be
explained with the available results.
- 13 -
Machining increased the hardness of the material with respect to the as-received
sample. This increase, however, was not associated with increase in recovered indentation
depth. This indicates that cold work must be followed by an annealing process to preserve the
ability to stress induce martensite.
Recovered indentation depth [nm]
320
700
300
280
200
heat-treated NiTi chips
260
800
400
240
600
as-received
500
as-received NiTi
300
as-machined NiTi
700C
as-machined
220
Linear (heat-treated NiTi chips)
200
180
160
100
110
120
130
140
150
160
170
180
190
200
Microhardness [HV]
Fig. 12: Superelasticity vs. microhardness
320
Recovered indentation depth [nm]
300
700
280
800
260
400
200
300
as-received
240
600
220
annealed chips
500
as received NiTi
as-machined
200
as machined
700C
180
Linear (annealed chips)
160
140
2.5
2.7
2.9
3.1
3.3
3.5
3.7
3.9
4.1
Nanohardness [GPa]
Fig. 13: Superelasticity vs. nanohardness
5. CONCLUSIONS
1. Recrystallization occurred after annealing for 1 hour at about 600ºC and was followed
by grain growth at 700 and 800ºC.
2. EDX analyses showed that precipitates increased up to about 500ºC then decreased
again above 600ºC due to solutionizing.
3. Microhardness readings were influenced by elastic recovery whereas nanohardness
readings were not affected by superelasticity. Microhardness was also more sensitive
to changes in the microstructure, e.g. presence of dislocation tangles and
recrystallization, than nanohardness.
4. Superelasticity of annealed cold-worked NiTi was influenced by arranged
dislocations and precipitation. Low temperature annealing had the advantage of
improving superelasticity while preserving hardness.
5. Generally, superelasticity of annealed cold-worked NiTi increased as micro- and
nanohardnesses increased.
- 14 -
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