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Micropyretic synthesis of NiTi in propagation mode
Acta Materialia 51 (2003) 2549–2568 www.actamat-journals.com
Micropyretic synthesis of NiTi in propagation mode G.K. Dey ∗ Materials Science Division, Bhabha Atomic Research Centre, Trombay, Mumbai 400 085, India Received 2 May 2002; accepted 16 December 2002
Abstract The microstructure of NiTi formed by micropyretic synthesis, carried out in the propagation mode, was investigated with preheat and Ti particle size as the two process variables. These variables were found to exert a significant influence on the microstructure. as well as the process. Unstable combustion was observed in specimens combusted without preheat. The microstructure of specimens undergoing unstable combustion showed some features which were similar to those developing during rapid solidification of NiTi in terms of the nature of the phases and their morphology. The microstructure essentially comprised the martensitic phase, the B2 parent phase and the Ti2Ni intermetallic phase. In some specimens the Ti11Ni14 phase was found to form in an unusual lamellar morphology. The microstructural features of the micropyretically synthesized alloy were compared with those of the conventionally processed alloy. The structures of the B2, Ti2Ni and Ti11Ni14 phases and of their interfaces with the matrix phase were examined by high resolution electron microscopy (HREM). The mechanism of synthesis of the alloy from elemental powders was studied by stopping the combustion front and carrying out detailed microstructural characterization around the arrested front. 2003 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Transmission electron microscopy; HREM; Combustion synthesis; NiTi; Shape memory
1. Introduction NiTi has emerged as a very important intermetallic compound because of its shape memory property and hence has been the focus of many studies [1–4]. The shape memory effect in this alloy and the phase transformation responsible for it have been studied in considerable detail by many workers [2–6]. The extent of the shape memory effect and the temperature range over which it is exhib-
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ited depend strongly on the composition of the alloy, and in order to realize the maximum benefits of this effect, it is essential to have an alloy of exact stoichiometry and very good homogeneity [2,4]. This alloy has been fabricated both by conventional casting and working route as well as by powder metallurgy techniques. Conventional melting practices, like induction or arc melting, do not very easily yield an alloy of the desired stoichiometry and homogeneity. In the powder metallurgy technique, consolidation of the pre-alloyed powders has been undertaken [7,8]. The advantage of this process is that the exact stoichiometry could be achieved rather easily. The mechanical properties of alloys made by this technique agree well
1359-6454/03/$30.00 2003 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. doi:10.1016/S1359-6454(03)00055-7
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with those exhibited by alloys prepared by other processing techniques [8,9]. Since the process of preparing prealloyed powders is rather cumbersome, there have been studies involving the use of elemental powders for the preparation of the alloys by vacuum hot pressing and by micropyretic synthesis [9–13]. Micropyretic synthesis, also known as combustion synthesis, is a method of synthesizing intermetallic compounds where a very homogeneous alloy of the desired stoichiometry can be obtained by a single step process [11–15]. Micropyretic synthesis as a means of materials processing, presents an attractive alternative to the conventional powder metallurgy technique [13] because of several special and attractive features, viz., self-generation of energy, high reaction temperatures (1500 to 4000 °C), short processing times (a few seconds to a few minutes), high temperature gradients, and high heating and cooling rates. Many intermetallic compounds, particularly the aluminides of Ni and Ti, are amenable to synthesis by micropyretics because of their high adiabatic temperatures and high heats of mixing [11,13,15]. The microstructures of the alloys synthesized by this process are very homogeneous, thus eliminating the need for homogenization heat treatments or the need for complex thermomechanical processing [16,17]. In many medical applications such as body implants, a porous NiTi alloy is preferred because it allows the tissues to grow around the pores, making the bonding between the implant and the tissues more secure. Microstructural studies on this material, produced by this process, show that it is possible to get a variety of microstructures and a variety of densities in this alloy by suitably choosing the processing condition [18]. It is therefore possible to tailor the microstructure of this alloy and the porosity for a variety of applications by choosing the right processing conditions. The transformation temperature, which is a function of the microstructure, can also be controlled by controlling the microstructure of the alloy. The NiTi alloy made by micropyretic synthesis has been found to have properties comparable to those exhibited by the alloy made by the conventional fabrication routes [11,13,15,19–22]. It has been shown by Yi and Moore that the synthesis of NiTi can occur in the thermal explosion mode [19].
It has also been shown that since the reaction is weak, preheating is needed to propagate the reaction [21]. It has been demonstrated by Zhu et al. [22] that explosively formed NiTi and conventionally processed NiTi have some differences; an important difference pertains to the presence of an FCC phase with a Ti8Ni stoichiometry in the former [22]. However, their study has been confined to specimens which were heat-treated after thermal explosion [22]. The need for a detailed analysis of the microstructure has been stressed in some earlier studies [19,20]. The earlier studies deal with the synthesis of the alloy in the explosive mode and not in the propagation mode of synthesis. The work reported in the present paper was undertaken with the following objectives: (1) To study the micropyretic synthesis process in the propagation mode with preheat and Ti powder size as variables. (2) To investigate the effect of these process parameters on the microstructure of the alloy. The microstructure was examined in the as-synthesized state with a view to establishing the conditions under which the alloy could be synthesized and used directly without any thermal or thermomechanical treatment. In this study Ni and Ti particles of very small sizes were used for the first time (Table 1). (3) To examine the origin of the unstable combustion front and its effect on the microstructure of the alloy. (4) To study the nature of the phases forming after various processing conditions in detail by diffraction and high resolution electron microscopy (HREM) and to compare the results with those pertinent to conventionally processed alloys. (5) To look into the origin of the diffuse intensity distribution in selected area electron diffraction patterns (SAED) obtained from B2 phase in this alloy, which transforms to the martensitic phase. (6) To study mechanism of synthesis by stopping the combustion front and carrying out a
Table 1 Purity and particle size of the powders used Element
Purity
Powder size
Ni Ti
99.7 99.5
3 µm ⫺150 µm, ⫺45 µm, ⫺20 µm
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detailed microstructural characterization around the stopped combustion front.
2. Experimental High purity Ni and Ti powders (purity and size indicated in Table 1) were weighed accurately and mixed in a Spex model 8000 ball mill for 1200 s without the addition of any liquid medium. Subsequently, the mixed powders were pressed into rectangular bar specimens in a double acting press at a pressure of 140 MPa. The pressed compacts were combusted in an inert argon atmosphere, as well as in air by igniting from one end. Some specimens were heated in a furnace before combustion. Differential scanning calorimetery (DSC) measurements were made in a Perkin Elmer DSC II unit. A Cambridge scanning electron microscope (SEM) equipped with a Princeton Gamma-Tech detector, was used for the examination of the microstructure. The energy dispersive spectroscopy (EDS) results were quantified by using a standardless quantification approach. The results of standardless quantification were quite close to those obtained by quantification done by using appropriate standards. Specimens for transmission electron microscopy (TEM) were prepared by dimple grinding followed by ion milling in a Gatan ion mill. The thin foils so obtained were examined in a JEOL 2000 FX electron microscope equipped with a Kevex detector. HREM was carried out in a JEOL 3010 microscope having a point to point resolution of 0.21 nm.
Fig. 1. A typical sample where the unstable combustion front could be seen.
smaller powder size (⫺20 µm). A typical sample where the unstable combustion front could be seen is shown in Fig. 1. It is to be noted that no preheating of the specimen was carried out in these cases. The product obtained when NiTi specimens were synthesized without preheat and using ⫺45 µm Ti powder showed oscillatory combustion. Though specimens made under identical conditions but using ⫺20 µm Ti powder also seemed to show oscillatory combustion, a careful examination of specimens (Fig. 2) revealed that the nature of combustion in these specimens could be better placed in the category of spin or spiralling type rather than
3. Results 3.1. Effect of particle size and preheat on the combustion process The Ti particle size alone was varied in this study because Ti is the higher melting component and the particle size of the higher melting element has a much stronger influence on the synthesis process. The combustion front was found to be unstable when ⫺45 µm Ti powder was used and continued to remain unstable with the use of a still
Fig. 2. SEM micrograph showing the cross section of the specimen undergoing spiraling combustion. The plane of the figure is perpendicular to the direction of propagation of the front.
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pure oscillatory combustion. The striations seen in Fig. 2 were along the direction of propagation of the front unlike the ones seen in Fig. 1 which were perpendicular to this direction. Though Ti powders larger than ⫺45 µm were also used in this study (see Table 1), the results discussed here pertain to the two aforementioned particle sizes because here the microstructure was found to depend strongly on the particle size. In instances where the Ti particle size was larger than 100 µm, unreacted Ti particle cores could be seen in the microstructure. The effect of preheating to a temperature of 573 K was examined in detail. Preheating to temperatures considerably higher than this led to a total loss in shape of the green compact due to extensive melting. 3.2. Phases and microstructures It was found that these changes in the combustion behavior were associated with significant changes in the microstructure of the alloy. This could possibly be attributed to local variations in the velocity of the combustion reaction and the variations in the post combustion solidification process. The microstructures of the following types of specimens were examined in considerable detail: (1) Specimens synthesized with a Ti particle size of ⫺20 µm and after preheating the green compact; (2) Specimens synthesized with a Ti particle size of ⫺20 µm but without preheating the green compact; (3) Specimens synthesized with a Ti particle size ⫺45 µm and after preheating the green compact; and (4) Specimens synthesized with a Ti particle of ⫺45 µm and without preheating the green compact. Typical DSC thermograms corresponding to specimens made with Ti powders of ⫺20 µm size and without preheat are shown in Fig. 3. The transformation temperatures, as determined from this thermogram (Ms = 335 K and As = 370 K) were higher than those of NiTi synthesized by other techniques [4]. It was felt that a comparison of the microstructures developed by unstable combustion with that produced by rapid solidification was pertinent because both these processes are rapid processes of intermetallic compound formation; in fact Li and Sekhar [23] have observed that in the case of
Fig. 3. A typical DSC thermogram obtained from a specimen synthesized without preheating and using ⫺20 µm Ti powder. The heating, as well as cooling, rate of 10 K/min was used.
unstable conditions of combustion synthesis, cooling rates of the order of 104 K/s may be generated. They have also shown that unstable combustion of NiAl produces a microstructure which has similarities with that produced by rapid solidification [23]. 3.2.1. Specimens synthesized with a Ti particle size of ⴚ20 mm and after preheating the green compact The microstructure of the specimens made under this processing condition essentially comprised equiaxed grains of the B2 phase. In addition to this there were regions, which showed a martensitic structure. Twins that were observed in this microstructure were ⬍011⬎ type II, {111} type I and {011} type I twins. Fig. 4a shows a ⬍011⬎ type II twin, whereas Fig. 4b and c show {111} type I and {011} type I twins, respectively. A HREM micrograph of the ⬍011⬎ type II twin is shown
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Fig. 4. (a), (b) and (c) are bright field micrographs showing ⬍011⬎ type II twin, {111} type I and {011} type I twins, respectively. (d) HREM micrograph of the ⬍011⬎ type II twin.
in Fig. 4d. The twin interfaces were flat and did not show any ledge or unevenness, but they showed a broad strain contrast. In addition to the martensitic phase, it was possible to see the presence of a second phase, which showed diffraction patterns very similar to those of the Ti4Ni2O phase (Fig. 5a). Since this phase is known to form during the solidification of the
alloy, no simple crystallographic orientation relationship would be expected to exist between this phase and the B2 phase, and this was verified in the present study (Fig. 5a). Fig. 5b is a bright field (BF) micrograph showing a particle of this phase. Fig. 5c shows a HREM micrograph of this phase and its interface with the B2 matrix. The interface is devoid of any steps or ledges. It could
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Fig. 5. (a) SAED pattern from the Ti4Ni2O phase. (b) Bright field micrograph from the Ti4Ni2O phase. (c) HREM micrograph showing the interface between the Ti4Ni2O phase and the matrix phase. (d) SAED form the B2 phase showing diffuse intensity indicative of the presence of the instability in the B2 phase. (e) HREM micrograph showing the 110 planes of the B2 phase.
be seen that the matrix lattice had not been imaged. This was because of the fact that no low index zone of the matrix was parallel to the interface in this imaging condition. Besides the martensitic regions, regions which essentially had the structure of the parent B2 phase were also seen. Fig. 5d shows a SAED pattern
from such a region. In addition to the spots due to the B2 Phase, it was possible to see diffuse intensity distribution or streaks which were indicative of the presence of some kind of instability in the B2 phase. Fig. 5e is an HREM image of the B2 phase corresponding to the [110] zone showing the (110) planes of this phase.
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Fig. 5.
3.2.2. Specimens synthesized with a Ti particle size of ⴚ20 mm and without preheating the green compact The nature of combustion in this case was of the spin type and not oscillatory. TEM examination of the specimens showed very fine grains (Fig. 6a), SAED patterns from which indicated that these were of the ordered B2 phase. It should be noted that the extremely small size of these grains could only have been produced by rapid solidification of the alloy. No martensitic phase could be noticed in the B2 grains. The microstructures at the centre of the band and at the edge were similar. In addition to the B2 phase, it was possible to see some ellipsoidal-shaped particles in this microstructure. One typical SAED pattern from such a particle is shown in Fig. 6b. It was possible to index this pattern and many others taken from these particles in terms of the Ti4Ni2O phase.
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Continued
3.2.3. Specimens synthesized with a Ti particle size of ⴚ45 mm and after preheating the green compact Fig. 7 shows a typical optical micrograph from such a specimen. It can be seen that the microstructure comprised dendritic grains, surrounded by a eutectic phase mixture at the grain boundaries. This microstructure was somewhat similar to that observed by Zhu et al. [22] in this alloy. A detailed TEM examination of this microstructure has been carried out and the results of these investigations will be presented elsewhere [24]. The eutectic phase mixture at the grain boundary consisted of the Ti4Ni2O and the B2 phases. Some pores could also be seen in the microstructure, which is an inherent feature of the microstructure produced by micropyretic synthesis. The pores were of two types: (1) Very small pores present at the grain boundaries as well as in the grain interiors; and
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Fig. 7. Optical micrograph depicting the microstructure of combusted specimens prepared with ⫺45 µm Ti particles and preheating of compact.
Fig. 6. (a) Bright field micrograph showing fine equiaxed grains in unpreheated specimens made with a Ti particle size of ⫺20 µm. (b) [110] SAED pattern from the Ti4Ni2O phase.
(2) Much larger pores present mostly at the grain boundaries. The grain size of these specimens was much larger than that of the other three types of specimens. 3.2.4. Specimens synthesized with a ⴚ45 mm Ti particle size without preheating the green compact The microstructure of specimens obtained by this type of combustion was found to contain more or less equiaxed grains of the B2 phase. SAED patterns from these grains indicated that the B2 phase was fully ordered. A disordered bcc phase has been observed in the case of NiTi synthesized by rapid
solidification [25]. When the equiatomic NiTi alloy is slowly cooled from temperatures of the order of 1273 K, the ordering reaction occurs at temperatures between 973 K to 873 K. Since this reaction is diffusion controlled, the effect of rapid solidification will be to partially suppress it. In this case no evidence for a partial suppression of order could be found. The examination of a large number of diffraction patterns form B2 regions showed diffuse intensity maxima indicative of the presence of an omega-like instability. HREM examination of the matrix B2 phase carried out along the [110] zone yielded results which were very similar to those obtained in the case of specimens synthesized without preheat and using ⫺20 µm Ti particles. The Ti4Ni2O phase also was present in these specimens (Fig. 8a). The faceting of the particles of this phase was much less evident in this case. Fig. 8b shows a BF micrograph of a matrensitic region in this microstructure. The ⬍011⬎ type II, {111} type I, and {011} type I twins could be observed in the martensitic regions. The scale of the martensitic microstructure was finer than that observed in the case of the conventionally produced alloy. HREM of the twin interfaces in the martensitic region showed a broad strain contrast.
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[111]Ti11Ni14 / / (111)M The lattice parameter of the phase Ti11Ni14, as determined by SAED, was found to be close to 0.730 nm. Simulated diffraction patterns with this value of the lattice parameter and corresponding to the SAED patterns in Fig. 9b and c are shown in shown Fig. 9d and e, respectively. Fig. 9f shows a complex diffraction pattern from a region which had at least two variants of this phase. The key to this diffraction pattern (Fig. 9g) indicates how this pattern had resulted from the two variants. Some of the spots in Fig. 9f are due to double diffraction. Diffraction along with dark field (DF) microscopy revealed that the parallel lamellae in any given region belonged to the same variant, whereas lamellae belonging to different variants were not parallel to each other. Fig. 9h shows a typical interface between the B2 phase and the Ti11Ni14 phase. The interface was not flat and had a wavy nature. Fig. 9i shows a HREM micrograph of a region containing two variants (marked as A and B) of this phase. 3.3. Microstructural characterization of the reaction front
Fig. 8. (a) Particles of the Ti4Ni2O in the matrix. (c) Bright field micrograph showing the martensitic phase.
These interfaces were flat and did not show any ledge or unevenness. In this regard the observations made here were very similar to those made in specimens made with ⫺20 µm Ti powder after preheating the green compact. Besides the martensitic regions, there were regions which had a lamellar microstructure (Fig. 9a). SAED patterns from such a region are shown in Fig. 9b and c. These patterns could be interpreted in terms of the B2 phase and the Ti11Ni14 phase (rhombohedral unit cell with a = 103.65). The orientation relationship between this phase and the B2 matrix was found to be: (3¯ 21)Ti11Ni14 / / (1¯ 10)M
The mechanism of synthesis was examined in the case of unpreheated specimens synthesized with ⫺45 µm Ti powder. The combustion front was frozen by using a technique described in Ref. [16]. In order to examine the nature of the reaction occurring at the front, the microstructure of the specimen was examined (1) away from the front on the powder side, (2) at the front and (3) on the intermetallic side of the front. Fig. 10a shows the frozen combustion front. Depending upon the extent of reaction, the specimen could be divided in to three distinct regions: unreacted region, partially reacted region and fully reacted region. A description of the microstructure of these three regions follows. 3.3.1. Unreacted region Fig. 10b shows the powder side of the front at a distance at which the microstructure was not affected by the heat generated at the front. The Ti and the Ni powders could be readily distinguished
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Fig. 9. (a) Bright field micrograph showing a lamellar microstructure. SAED pattern from the lamellar region, (b) zone axis [111] B2, and (c) zone axis [1¯ 12] B2. (d) Simulated and superimposed SAED patterns. Open circles represent B2 reflections and filled circles represent Ti11Ni14 reflections. (e) Simulated and superimposed SAED patterns. Black circles represent B2 reflections and open circles represent Ti11Ni14 reflections. (f) Complex SAED pattern showing diffraction spots due to B2 phase, two variants of the Ti11Ni14 and spots due to double diffraction. (g) Simulated and superimposed SAED patterns of B2 and Ti11Ni14 phases. Black circles represent B2 reflections, open circles and filled circles represent reflections of two variants of Ti11Ni14. (h) HREM micrograph showing the interface between the B2 and the Ti11Ni14 phase. (i) HREM micrograph showing the B2 and two variants of the Ti11Ni14 phase (designated A and B).
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Fig. 9.
by comparing the secondary electron image with X-ray mapping. The Ti powder particles were all surrounded by Ni particles and instances of a Ni particle touching another Ni particle were very few. Ni was the lower melting ingredient and therefore was likely to melt first and surround the particles of Ti during the initial stages of synthesis. The Ti-Ni reaction is known to have high adiabatic temperature and enthalpy of formation and a low ignition temperature [10]. A study investigating the mechanism of synthesis in Ti-C-Ni-Mo alloys has shown that the Ti
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Continued
particles undergo a phase transformation wherein the a phase transforms to the b phase and then on cooling transforms to the a phase martensitically [26]. In the present study a close examination of Ti particles close to the reaction front revealed that a transformation of this nature did occur. This is clearly shown in Fig. 10c, where it could be seen that many of the Ti particles contained the martensitic phase in them. This observation was indicative of the fact that in these regions the temperature had increased beyond the a→b transition temperature (1136 K). The presence of the mar-
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Fig. 9.
tensitic phase could not be seen in particles in regions further removed from the reaction front because these had not been heated sufficiently for their temperature to have crossed the a→b transition temperature. In this description even the region in which martensitic transformation had occurred in the Ti particles has been categorized as a part of the unreacted region because the particles had not reacted with each other to any appreciable extent. 3.3.2. Partially reacted region The stopped combustion front comprised the region where the final microstructure had formed (Fig. 10a). One side of it was a region where Ni had melted and reacted with Ti either fully or partially. The melting of Ni occurred at some distance ahead of the front. This region extended from the place where the final microstructure could be seen to the point where the Ni had just melted without undergoing any substantial reaction. On moving
Continued
further towards the fully reacted region it was observed that Ni melted first and enveloped the Ti particles (Fig. 10d), leading to the formation of a Ni rich intermetallic phase which had a stoichiometry very close to Ni4Ti as revealed by EDS. An EDS profile depicting this is shown in Fig. 10e. The full scale reading in this and the other spectra shown in this work is 5000 counts. In addition, the formation of an intermetallic containing a higher amount of Ti and having a stoichiometry close to Ni2Ti could be seen (Fig. 10f). The regions from where the EDS spectra shown in Fig. 10e and f have been taken are marked as regions A and B in Fig. 10d. EDS analysis also showed that there were regions where the large Ti particles had an unreacted core (region marked as C in Fig. 10d). On moving still further towards the fully reacted region, it was possible to see two types of regions which had stoichiometries close to Ti3Ni and NiTi, respectively. These regions are marked in the micrograph shown in Fig. 10g. None of the Ti par-
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Fig. 10. (a) Micrograph showing the frozen combustion front. (b) Micrograph showing the powder side of the front at a distance at which the microstructure is not affected by the heat generated at the front. (c) Micrograph showing Ti and Ni particles. (d) Micrograph showing the melting of Ni which envelops the Ti particles and leads to the formation of a Ni rich intermetallic phase, which has a stoichiometry very close to Ni4Ti. (e) EDS profile depicting the formation of phase having stoichiometry close to Ni4Ti. (f) EDS profile showing the formation of intermetallics containing higher amounts of Ti and having stoichiometry close to Ni2Ti. The points from where EDS profiles (e) and (f) have been taken are marked as A and B in the micrograph (d). Point C in this micrograph shows an unreacted Ti core. (g) Micrograph showing an area very close to the fully reacted region. Points from where profiles showing stochiometry of Ti3Ni and close to NiTi are have been taken are marked as D and E in this micrograph.
could be construed. The molten Ni dissolved the smaller Ti particles to form Ni rich intermetallic compounds, such as Ni4Ti. The larger particles of Ti were not consumed quickly. Instead, these formed intermetallics such as Ti3Ni2 at the rim and Ti4Ni at the centre of the particle. The Ni-rich intermetallics then reacted with more of Ti and Ti rich intermetallics to form the equiatomic NiTi phase. It should be noted that in addition to the NiTi phase, the formation of the Ti2Ni or the Ti4Ni2O phases also occurred. The presence of this phase could be seen at the grain boundaries in the fully reacted side of the specimen.
ticles showed the formation of the martensitic phase. presumably because the Ti particles had undergone alloying and therefore the a→b transition temperature corresponding to these regions had changed substantially. Based on the observations made on the partially reacted region, the following sequence of microstructure formation
3.3.3. Fully reacted region The microstructure of this region was very similar to that of the fully reacted specimen in which the combustion front had not been stopped. Here a two phase microstructure could be seen (Fig. 11a). Typical EDS profiles from the center of the grain and the grain boundary, shown in Fig. 11b and c, respectively, indicated that while the composition of the center of the grain was close to NiTi, that of the grain boundaries corresponded to Ti2Ni. The size of the NiTi grains close to the front was found to be smaller than that of those away from the front.
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Fig. 10. Continued
4. Discussion 4.1. Effect of particle size and preheat on the process of synthesis It has been observed that the particle size influences the combustion process by affecting the combustion velocity [10]. The combustion velocity remains constant over a range of small particle sizes. An increase in the particle size beyond this range leads to a drastic reduction in the velocity of propagation, followed by a region where the decrease in the velocity is very weakly related to the particle size [10]. In most of the earlier studies, the Ni particle size has been altered [13]. The generation of heat by the combustion reaction and the rate of heat loss from the reaction front govern the rates of propagation, wave stability and the maximum combustion temperature during micropyretic synthesis [13]. Heat is dissipated away from the reaction front by heat transfer to the
adjacent reactant mixture, the temperature of which is below the ignition temperature, and by heat loss to the environment. A change in the reaction velocity and stability of the front will occur if these factors undergo any perturbation. Highly exothermic reactions are self-sustaining and as the enthalpy of reaction decreases, a limit is reached at which the combustion is extinguished. A nonsteady state of the reaction front results if the perturbations persist. The many kinetic reasons for a departure from steady-state propagation include particle size, compaction pressure and presence of diluents. Instability of the combustion front can arise from either inadequate heat generation during the combustion process or from a very high density of the green compact. Whereas the former leads to an instability of the front due to the non-availability of sufficient heat to sustain it, the latter causes instability because the high conductivity of the green compact leads to a rapid dissipation of heat, leading to a thermal imbalance. In the present
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Fig. 11. (a) Micrograph showing a two phase microstructure in the fully reacted region. (b) and (c) EDS profiles from the center of the grain and the grain boundary respectively.
study, the combustion process was found to be unstable when ⫺45 µm and ⫺20 µm Ti powders were used. The instability of the combustion process seemed to be due to the high density of the green compact. On preheating the specimen, however, the front could be made stable. This was indicative of the fact that the instability of the combustion front was also due to the inadequacy of the heat produced during the combustion process. The microstructure generated during the synthesis process was dependent on the Ti particle size. It will be shown later that the Ti particle size had an influence on the grain morphology and the nature of the phases forming. Microstructural studies on unpreheated specimens made from two Ti particle sizes indicated some important differences. This was because of the fact that when the particle size was ⫺45 µm the
nature of the unstable combustion corresponded to the oscillatory type propagation whereas when the particle size was reduced further to ⫺20 µm the nature of the combustion was found to change from the oscillatory to the spin combustion mode [13]. The effect of the initial temperature of the specimen on the process of combustion has been investigated by many workers [13]. The combustion temperature changes with the initial temperature of the specimen. In the case of NiAl, it has been shown by Li and Sekhar that the adiabatic temperature remains constant with the only change being an increase in the molar melted fraction of the product phase from 42 to 100% [27]. In the present study, the combustion front was found to have become stable upon preheats of the specimen and its microstructure was found to undergo a transition from equiaxed to dendritic with a substantial variation
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in the internal structure of the grains. The effect of preheat appeared to be the same as seen with other intermetallic compounds [16,17,28]. 4.2. Effect of process variables on the microstructure The microstructure was found to be considerably influenced by the two process variables, namely, preheat and Ti particle size and the effects of these variables were noticed on the morphology of the grains, the nature of the phases present and the scale of the microstructure. The important differences in the overall grain morphology and phase content are summarized in Table 2. The observation of a very fine grained microstructure in the case of specimens synthesized without preheat and with ⫺20 µm Ti powder was indicative of a high cooling rate. The grain size and the substructure were finer than those in the specimens described in Section 3.2.4, indicating that the spin type of combustion led to a higher Table 2 Effect of process variables on the overall microstructure
cooling rate than that encountered in the oscillatory mode of combustion of the alloy. The grain morphology transformed to a dendritic type on preheating of specimens synthesized with ⫺45 µm Ti. Dey and Sekhar [16] have examined the effect of preheating on the microstructure of NiAl alloys and have shown that the extent of preheat influences not only the morphology of the grains, but also the process of solute segregation. They have shown that in binary NiAl, preheating transforms the microstructure from the equiaxed to the dendritic. In NiAl alloys, containing small alloying additions such as Fe, Cr and V, it has been observed that preheating prior to synthesis results in the formation of second phase particles at the grain boundaries [16]. Such a change in microstructure and segregation of alloying elements to the grain boundaries occurs in the preheated specimens because in these the amount of the melted fraction is high and the specimens cool slowly after the combustion process, leading to solidification at a slow rate. In the case of the NiTi alloy, the pre-
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heating of the specimens had lead to the formation of dendritic grains. Moreover, segregation of the elements had occurred, facilitating formation of the eutectic microstructure at the grain boundaries. A comparison was made between this microstructure and that produced by conventional solidification of NiTi ingots and a remarkable similarity was noticed because of the fact that in either case the alloy was heated well above its melting temperature and then allowed to cool slowly. In specimens with a Ti particle size of ⫺20 µm, the grains were equiaxed even after preheating. This was so because heat dissipation occurred at a more rapid rate as compared to the case of the specimens with a powder size of ⫺45 µm. It was interesting to note that the B2 phase showed almost perfect order even in specimens, which had been synthesized with an unstable front. Though this phase formed by a rapid processing technique involving the reaction of the elemental powders of Ti and Ni, the mixing of the elements was quite homogeneous. The SAED patterns from the B2 phase showed streaks (Fig. 5d). Such streaks have also been noticed in the B2 phase of the NiTi alloy produced by the conventional route [2] and are quite similar to those observed in diffraction patterns obtained from b phase of the w forming systems [29]. The b→w transformation has been mostly found to occur in the group IV metals [30]. Experimental work as well as theoretical studies have been carried out to explain the diffuse w scattering [30]. HREM studies have been performed by Schryvers and Tanner [29] to ascertain the changes occurring in the b lattice when the diffuse w scattering initiates. Their HREM studies have shown that the diffuse w scattering is due to the presence of short strings of one and two-dimensional w phase. HREM was carried out in the present study along the [110] B2 zone axis. It could be seen that the HREM image (Fig. 5e) did not show the presence of features like the ω strings observed by Schryvers and Tanner [29]. The diffuse maxima seen in the diffraction pattern, therefore, could not be ascribed to the type of structural transformation observed in the case of w forming alloys. In the case of NiTi, the formation of an ω type phase after aging of the B2 phase has not been reported. A detailed HREM study of the diffuse
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intensity distribution in the diffraction patterns from the B2 phase in NiTi has not been carried out even in a conventionally produced NiTi alloy. Moine et al. [31,32] have examined the premartensitic diffraction effects which occur just above the martensitic transition temperature by examining the electron diffraction patterns. They have attributed the observed extra reflections (not diffuse intensity maxima) to the occurrence of two types of particles and the presence of lattice displacement waves in these particles. In this study, however, no extra reflections in the SAED patterns from the B2 phase could be noticed. The structure of the martensitic phase observed in this study was very similar to that produced by conventional processing. This could be ascertained both by conventional microscopy, as well as by HREM. HREM examination of the twin interfaces indicated that these were flat and did not show any ledges or unevenness. The observations made in this study on the structure of the interface were similar to those made by Nishida et al. [5] and were at variance with those made by Knowles [1]. The scale of the microstructure, however, was much finer than that in the case of conventional processing. The Ms temperature of the specimens examined in this study was found to be higher than that of the conventionally produced NiTi alloy. It has been mentioned earlier that the microstructure produced under conditions of unstable combustion were found to have features, such as very fine grain size, which suggested the occurrence of solidification under a high cooling rate. Therefore it is worth comparing the Ms of specimens of NiTi synthesized by unstable combustion with those produced by rapid solidification. Igaharo and Wood [25] have shown that the Ms temperature of NiTi is lowered by rapid solidification. Factors which could affect the transformation behavior of this alloy following rapid solidification include the high concentration of quenched-in vacancies, the dislocation density, the degree of disorder and the grain size. Since the equilibrium concentration of vacancies increases with temperature in an exponential manner, it is expected that a rapidly solidified alloy will have a higher concentration of vacancies as compared with an alloy quenched from the solid
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state [25]. The vacancies formed in excess tend to form vacancy clusters, which may collapse to yield dislocation loops. It has also been seen that some of these vacancies anneal out if the cooling rate during the post solidification stage is not high enough [25]. The strain field associated with the dislocations and dislocation loops tend to promote martensite formation and consequently raise the Ms by providing a favorable interaction with the strain fields of the martensite nuclei [25]. The other factors affecting the Ms value are the grain size, the extent of order in the parent phase and the amount of oxygen present in the material. In this study it was observed that the Ms temperature for the alloy synthesized under an unstable condition was higher than that for an alloy produced by conventional solidification. This was indicative of the fact that the aforementioned factors had indeed played an important role in relation to the Ms temperature for specimens examined in this study. It has been shown by Moore and Yi [33] that a higher cooling rate from the molten state leads to a dendritic morphology of the Ti4Ni2O phase (designated in their study as the Ti2Ni phase), whereas a slow cooling rate leads to a faceted morphology of this phase. In the present study it was seen that the morphology of the Ti4Ni2O phase indeed depended on the processing condition. Besides the eutectic morphology a faceted morphology was also observed in specimens produced under a different set of processing conditions. The Ti4Ni2O phase has a fcc structure of the Fe3W3C type with a lattice parameter a = 1.13193 nm, while the Ti2Ni phase also has an identical structure with a lattice parameter a = 1.13279 nm [34]. Ti2Ni can transform to a quasicrystalline structure on rapid solidification [35]. The quaiscrystalline phase forming on rapid solidification in a (Ti1-x Vx)2Ni alloy, specially after heat treatments at 673 K, shows a kind of local or short-range transnational order as evidenced by a more or less commensurate arrangement of electron diffraction spots along one or several two-fold directions [35]. Though the cooling rate during unstable synthesis was high, it was not high enough to lead to the formation of a quasicrystalline structure in this study. The Ti11Ni14 phase could be seen in one of the specimens in an unusual plate shaped morphology.
It may be noted that this phase has been found to form in the NiTi alloy after ageing at temperatures below 973 K and has been seen to adopt an oval plate morphology with lenticular cross-section [34]. Although the crystallographic orientation relationship between this phase and the B2 matrix observed in this study was the same as that seen earlier [34] in other studies [34], the lattice parameter of the unit cell was found to be larger (a = 0.720 nm) as compared to a = 0.661 nm reported in Ref. [34]. The formation of this phase in a lamellar arrangement, noticed in the present study has not been reported so far. This phase plays an important role in the development of the all round shape memory effect in the NiTi alloy [2]. Even though the morphology observed in this study resembled that of a martensitic phase, the interface structure and the substructure showing neither a high density of dislocations or of twins appeared to discount this possibility. The features of the Ti11Ni14 precipitates were suggestive of their having been formed by a diffusional transformation. This phase could be seen only in specimens synthesized with ⫺45 µm Ti. 4.3. Mechanism of synthesis The examination of the mechanism of synthesis could help in establishing the sequence of events leading to the formation of the final microstructure and in identifying those process parameters which control the microstructure. In this study the mechanism of synthesis was studied by stopping the combustion front, by making use of an arrangement of copper plates around the specimen, as described in Ref. [16]. Microstructural examination was carried out at different regions of the specimen after stopping the combustion front. Based on the microstructural examination of the three regions in the partially combusted specimen, the following sequence for the formation of the final microstructure from the powder particles could be proposed. The formation of the NiTi phase is triggered by the melting of the Ni particles, which then engulf the Ti particles. At first, Ni rich phases such as Ni4Ti and metastable phases form. The Ni rich phases then react with more of Ti and yield phases which are progressively richer
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in Ti. Ultimately, as the temperature increases, Ni reacts with Ti2Ni3 and Ti to form stoichiometric TiNi. Some amount of the Ti2Ni phase remains unreacted and becomes a part of the final microstructure. The mechanism of synthesis described here is pertinent to one set of processing conditions (Ti particle size and preheat). The mechanism may differ in minute details when examined for another set of processing conditions because the final microstructure depends on these conditions. However, the basic steps in the mechanism could be expected to remain the same.
5. Conclusions On the basis of the observations made in this study, the following conclusions could be drawn: 1. It was possible to generate a variety of microstructures in NiTi synthesized by micropyretic synthesis by appropriately choosing the processing conditions. These microstructures could range from the dendritic microstructures commonly encountered in ingots in the as-cast condition to equiaxed microstructures obtained after thermomechanical processing. Very fine grained microstructures could be obtained in regions undergoing unstable combustion. 2. The combustion process was unstable when preheating was not used. This instability appeared to be due to the fact that very fine Ti and Ni powders were used and the thermal conductivity of the resulting compacts was very good, promoting rapid heat dissipation. The nature of the instability was found to change from the oscillatory mode of combustion to spin type combustion with a decrease in the Ti particle size. This was presumably due to a progressive increase in the thermal conductivity of the compact with decreasing Ti particle size. 3. HREM examination of the B2 phase indicated that the instability in this phase, manifested in the form of diffuse maxima in SAED patterns, was not due to the formation of an w-like structure. 4. The nature of the martensitic phase was the same as that encountered in the conventionally
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processed alloy except for the fact that the size of the martensite plates was much finer. The martensite twin boundaries were devoid of any ledges or steps and were very similar to those observed in the conventionally produced alloy. 5. The Ti4Ni2O phase had the same structure and morphology as observed in the case of the conventionally produced alloy. The morphology of the phase was found to depend on the processing conditions. The Ti11Ni14 phase was found to show an unusual lamellar morphology. 6. The mechanism of synthesis of the alloy involved melting of Ni, followed by reaction with Ti to form Ni rich intermetallic compounds. These then reacted with more of Ni and Ti to give rise to intermetallic compounds with higher and higher amounts of Ti, finally leading to the formation of the NiTi phase.
Acknowledgements The author would like to thank Dr. K. Madangopal, Dr. D. Srivastava, Dr. P. Mukhopadhyay and Dr. P. K. De for many helpful discussions. Thanks are also due to Prof. J. A. Sekhar for his constant encouragement during the course of this work.
References [1] Knowles KM. Phil Mag 1982;A45:357. [2] Saburi T. In: Otsuka K, Wayman CM, editors. Shape memory alloys. Cambridge: Cambridge University Press; 1998. p. 49. [3] Knowles KM, Smith DA. Acta Metall 1981;29:101. [4] Perkins J, editor. Shape memory effect in alloys. NY: Plenum Press; 1975. [5] Nishida M, Ohigi H, Itai I, Chiba A, Yamauchi R. Acta Metall Mater 1995;43:1219. [6] Kudoh Y, Tokonami M, Miyazaki S, Otsuka K. Acta Metall 1985;33:2049. [7] Johnson A, Dominque JA, Reichem SH. J D Physique 1982;43(C4):285. [8] Johnson A, Dominque JA, Reichem SH. J D Physique 1982;43(C4):291. [9] Igaharo M, Wood JV. Powder Metall 1985;28:131. [10] Sckiguchi Y, Funami K, Funakubo H, Suzuki Y. J D Physique 1982;43(C4):279. [11] Moore JJ, Feng HJ. Progress in Materials Science 1995;39:243.
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[12] Munir ZA. Am Ceram Soc Bull 1988;67:342. [13] Munir ZA, Anselmi-Tamburini U. Mater Sci Reports 1989;3:277. [14] Merzhanov AG. In: Munir ZA, Holt JB, editors. Combustion and plasma synthesis of high-temperature materials. New York: VCH Publishers; 1990. p. 1. [15] Moore JJ, Feng HJ. Progress in Materials Science 1995;39:275. [16] Dey GK, Sekhar JA. Metall and Mater Trans 1997;28B:905. [17] Dey GK, Sekhar JA. Metall and Mater Trans 1999;30B:171. [18] Li BY, Rong LJ, Li YY, Gjunter VJ. Acta Mat 2000;48:3895. [19] Yi HC, Moore JJ. J Mater Sci 1989;24:3449. [20] Yi HC, Moore JJ. J Mater Sci 1989;24:3456. [21] Itin VI, Bratchikov AD, Dronin VN, Pribytkov GA. Sov Phys J 1981;24:1134. [22] Zhu M, Li TC, Liu JT, Yang DZ. Acta Metall Mater 1991;39:1481. [23] Li HP, Sekhar JA. J Mater Res 1993;8:2515.
[24] Dey GK. Unpublished work. [25] Igharo M, Wood JV. In: Froes FH, Savage SJ, editors. Processing of structural metals by rapid solidification. A S M International; 1987. p. 217. [26] Lasalvia JC, Kim DK, Lipsett RA, Meyers MA. Met and Mat Trans 1995;26A:3001. [27] Li HP, Sekhar JA. J Mater Sci 1995;30:4628. [28] Dey GK, Arya A, Sekhar JA. J Mat Res 2000;15:63. [29] Schryvers D, Tanner LE. In: Muddle BC, editor. Materials science forum. 1990;56-58:329. [30] Sinkler W, Luzzi DE. Acta Met Mat 1994;42:1249. [31] Moine P, Michal GM, Sinclair R. Acta Metall 1982;30:109. [32] Moine P, Michal GM, Sinclair R. Acta Metall 1982;30:125. [33] Moore JJ, Yi HC. Materials Science Forum 1990;5658:637. [34] Nishida M, Wayman CM, Honma T. Met Trans 1986;17A:1505. [35] Zhang Z, Kuo KH. J Microscopy 1987;146:313.
Micropyretic synthesis of NiTi in propagation mode G.K. Dey ∗ Materials Science Division, Bhabha Atomic Research Centre, Trombay, Mumbai 400 085, India Received 2 May 2002; accepted 16 December 2002
Abstract The microstructure of NiTi formed by micropyretic synthesis, carried out in the propagation mode, was investigated with preheat and Ti particle size as the two process variables. These variables were found to exert a significant influence on the microstructure. as well as the process. Unstable combustion was observed in specimens combusted without preheat. The microstructure of specimens undergoing unstable combustion showed some features which were similar to those developing during rapid solidification of NiTi in terms of the nature of the phases and their morphology. The microstructure essentially comprised the martensitic phase, the B2 parent phase and the Ti2Ni intermetallic phase. In some specimens the Ti11Ni14 phase was found to form in an unusual lamellar morphology. The microstructural features of the micropyretically synthesized alloy were compared with those of the conventionally processed alloy. The structures of the B2, Ti2Ni and Ti11Ni14 phases and of their interfaces with the matrix phase were examined by high resolution electron microscopy (HREM). The mechanism of synthesis of the alloy from elemental powders was studied by stopping the combustion front and carrying out detailed microstructural characterization around the arrested front. 2003 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Transmission electron microscopy; HREM; Combustion synthesis; NiTi; Shape memory
1. Introduction NiTi has emerged as a very important intermetallic compound because of its shape memory property and hence has been the focus of many studies [1–4]. The shape memory effect in this alloy and the phase transformation responsible for it have been studied in considerable detail by many workers [2–6]. The extent of the shape memory effect and the temperature range over which it is exhib-
∗
Tel.: 91-22-25595063; fax: 91-22-25505151, 91-2225505239. E-mail address: [email protected] (G.K. Dey).
ited depend strongly on the composition of the alloy, and in order to realize the maximum benefits of this effect, it is essential to have an alloy of exact stoichiometry and very good homogeneity [2,4]. This alloy has been fabricated both by conventional casting and working route as well as by powder metallurgy techniques. Conventional melting practices, like induction or arc melting, do not very easily yield an alloy of the desired stoichiometry and homogeneity. In the powder metallurgy technique, consolidation of the pre-alloyed powders has been undertaken [7,8]. The advantage of this process is that the exact stoichiometry could be achieved rather easily. The mechanical properties of alloys made by this technique agree well
1359-6454/03/$30.00 2003 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. doi:10.1016/S1359-6454(03)00055-7
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with those exhibited by alloys prepared by other processing techniques [8,9]. Since the process of preparing prealloyed powders is rather cumbersome, there have been studies involving the use of elemental powders for the preparation of the alloys by vacuum hot pressing and by micropyretic synthesis [9–13]. Micropyretic synthesis, also known as combustion synthesis, is a method of synthesizing intermetallic compounds where a very homogeneous alloy of the desired stoichiometry can be obtained by a single step process [11–15]. Micropyretic synthesis as a means of materials processing, presents an attractive alternative to the conventional powder metallurgy technique [13] because of several special and attractive features, viz., self-generation of energy, high reaction temperatures (1500 to 4000 °C), short processing times (a few seconds to a few minutes), high temperature gradients, and high heating and cooling rates. Many intermetallic compounds, particularly the aluminides of Ni and Ti, are amenable to synthesis by micropyretics because of their high adiabatic temperatures and high heats of mixing [11,13,15]. The microstructures of the alloys synthesized by this process are very homogeneous, thus eliminating the need for homogenization heat treatments or the need for complex thermomechanical processing [16,17]. In many medical applications such as body implants, a porous NiTi alloy is preferred because it allows the tissues to grow around the pores, making the bonding between the implant and the tissues more secure. Microstructural studies on this material, produced by this process, show that it is possible to get a variety of microstructures and a variety of densities in this alloy by suitably choosing the processing condition [18]. It is therefore possible to tailor the microstructure of this alloy and the porosity for a variety of applications by choosing the right processing conditions. The transformation temperature, which is a function of the microstructure, can also be controlled by controlling the microstructure of the alloy. The NiTi alloy made by micropyretic synthesis has been found to have properties comparable to those exhibited by the alloy made by the conventional fabrication routes [11,13,15,19–22]. It has been shown by Yi and Moore that the synthesis of NiTi can occur in the thermal explosion mode [19].
It has also been shown that since the reaction is weak, preheating is needed to propagate the reaction [21]. It has been demonstrated by Zhu et al. [22] that explosively formed NiTi and conventionally processed NiTi have some differences; an important difference pertains to the presence of an FCC phase with a Ti8Ni stoichiometry in the former [22]. However, their study has been confined to specimens which were heat-treated after thermal explosion [22]. The need for a detailed analysis of the microstructure has been stressed in some earlier studies [19,20]. The earlier studies deal with the synthesis of the alloy in the explosive mode and not in the propagation mode of synthesis. The work reported in the present paper was undertaken with the following objectives: (1) To study the micropyretic synthesis process in the propagation mode with preheat and Ti powder size as variables. (2) To investigate the effect of these process parameters on the microstructure of the alloy. The microstructure was examined in the as-synthesized state with a view to establishing the conditions under which the alloy could be synthesized and used directly without any thermal or thermomechanical treatment. In this study Ni and Ti particles of very small sizes were used for the first time (Table 1). (3) To examine the origin of the unstable combustion front and its effect on the microstructure of the alloy. (4) To study the nature of the phases forming after various processing conditions in detail by diffraction and high resolution electron microscopy (HREM) and to compare the results with those pertinent to conventionally processed alloys. (5) To look into the origin of the diffuse intensity distribution in selected area electron diffraction patterns (SAED) obtained from B2 phase in this alloy, which transforms to the martensitic phase. (6) To study mechanism of synthesis by stopping the combustion front and carrying out a
Table 1 Purity and particle size of the powders used Element
Purity
Powder size
Ni Ti
99.7 99.5
3 µm ⫺150 µm, ⫺45 µm, ⫺20 µm
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detailed microstructural characterization around the stopped combustion front.
2. Experimental High purity Ni and Ti powders (purity and size indicated in Table 1) were weighed accurately and mixed in a Spex model 8000 ball mill for 1200 s without the addition of any liquid medium. Subsequently, the mixed powders were pressed into rectangular bar specimens in a double acting press at a pressure of 140 MPa. The pressed compacts were combusted in an inert argon atmosphere, as well as in air by igniting from one end. Some specimens were heated in a furnace before combustion. Differential scanning calorimetery (DSC) measurements were made in a Perkin Elmer DSC II unit. A Cambridge scanning electron microscope (SEM) equipped with a Princeton Gamma-Tech detector, was used for the examination of the microstructure. The energy dispersive spectroscopy (EDS) results were quantified by using a standardless quantification approach. The results of standardless quantification were quite close to those obtained by quantification done by using appropriate standards. Specimens for transmission electron microscopy (TEM) were prepared by dimple grinding followed by ion milling in a Gatan ion mill. The thin foils so obtained were examined in a JEOL 2000 FX electron microscope equipped with a Kevex detector. HREM was carried out in a JEOL 3010 microscope having a point to point resolution of 0.21 nm.
Fig. 1. A typical sample where the unstable combustion front could be seen.
smaller powder size (⫺20 µm). A typical sample where the unstable combustion front could be seen is shown in Fig. 1. It is to be noted that no preheating of the specimen was carried out in these cases. The product obtained when NiTi specimens were synthesized without preheat and using ⫺45 µm Ti powder showed oscillatory combustion. Though specimens made under identical conditions but using ⫺20 µm Ti powder also seemed to show oscillatory combustion, a careful examination of specimens (Fig. 2) revealed that the nature of combustion in these specimens could be better placed in the category of spin or spiralling type rather than
3. Results 3.1. Effect of particle size and preheat on the combustion process The Ti particle size alone was varied in this study because Ti is the higher melting component and the particle size of the higher melting element has a much stronger influence on the synthesis process. The combustion front was found to be unstable when ⫺45 µm Ti powder was used and continued to remain unstable with the use of a still
Fig. 2. SEM micrograph showing the cross section of the specimen undergoing spiraling combustion. The plane of the figure is perpendicular to the direction of propagation of the front.
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pure oscillatory combustion. The striations seen in Fig. 2 were along the direction of propagation of the front unlike the ones seen in Fig. 1 which were perpendicular to this direction. Though Ti powders larger than ⫺45 µm were also used in this study (see Table 1), the results discussed here pertain to the two aforementioned particle sizes because here the microstructure was found to depend strongly on the particle size. In instances where the Ti particle size was larger than 100 µm, unreacted Ti particle cores could be seen in the microstructure. The effect of preheating to a temperature of 573 K was examined in detail. Preheating to temperatures considerably higher than this led to a total loss in shape of the green compact due to extensive melting. 3.2. Phases and microstructures It was found that these changes in the combustion behavior were associated with significant changes in the microstructure of the alloy. This could possibly be attributed to local variations in the velocity of the combustion reaction and the variations in the post combustion solidification process. The microstructures of the following types of specimens were examined in considerable detail: (1) Specimens synthesized with a Ti particle size of ⫺20 µm and after preheating the green compact; (2) Specimens synthesized with a Ti particle size of ⫺20 µm but without preheating the green compact; (3) Specimens synthesized with a Ti particle size ⫺45 µm and after preheating the green compact; and (4) Specimens synthesized with a Ti particle of ⫺45 µm and without preheating the green compact. Typical DSC thermograms corresponding to specimens made with Ti powders of ⫺20 µm size and without preheat are shown in Fig. 3. The transformation temperatures, as determined from this thermogram (Ms = 335 K and As = 370 K) were higher than those of NiTi synthesized by other techniques [4]. It was felt that a comparison of the microstructures developed by unstable combustion with that produced by rapid solidification was pertinent because both these processes are rapid processes of intermetallic compound formation; in fact Li and Sekhar [23] have observed that in the case of
Fig. 3. A typical DSC thermogram obtained from a specimen synthesized without preheating and using ⫺20 µm Ti powder. The heating, as well as cooling, rate of 10 K/min was used.
unstable conditions of combustion synthesis, cooling rates of the order of 104 K/s may be generated. They have also shown that unstable combustion of NiAl produces a microstructure which has similarities with that produced by rapid solidification [23]. 3.2.1. Specimens synthesized with a Ti particle size of ⴚ20 mm and after preheating the green compact The microstructure of the specimens made under this processing condition essentially comprised equiaxed grains of the B2 phase. In addition to this there were regions, which showed a martensitic structure. Twins that were observed in this microstructure were ⬍011⬎ type II, {111} type I and {011} type I twins. Fig. 4a shows a ⬍011⬎ type II twin, whereas Fig. 4b and c show {111} type I and {011} type I twins, respectively. A HREM micrograph of the ⬍011⬎ type II twin is shown
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Fig. 4. (a), (b) and (c) are bright field micrographs showing ⬍011⬎ type II twin, {111} type I and {011} type I twins, respectively. (d) HREM micrograph of the ⬍011⬎ type II twin.
in Fig. 4d. The twin interfaces were flat and did not show any ledge or unevenness, but they showed a broad strain contrast. In addition to the martensitic phase, it was possible to see the presence of a second phase, which showed diffraction patterns very similar to those of the Ti4Ni2O phase (Fig. 5a). Since this phase is known to form during the solidification of the
alloy, no simple crystallographic orientation relationship would be expected to exist between this phase and the B2 phase, and this was verified in the present study (Fig. 5a). Fig. 5b is a bright field (BF) micrograph showing a particle of this phase. Fig. 5c shows a HREM micrograph of this phase and its interface with the B2 matrix. The interface is devoid of any steps or ledges. It could
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Fig. 5. (a) SAED pattern from the Ti4Ni2O phase. (b) Bright field micrograph from the Ti4Ni2O phase. (c) HREM micrograph showing the interface between the Ti4Ni2O phase and the matrix phase. (d) SAED form the B2 phase showing diffuse intensity indicative of the presence of the instability in the B2 phase. (e) HREM micrograph showing the 110 planes of the B2 phase.
be seen that the matrix lattice had not been imaged. This was because of the fact that no low index zone of the matrix was parallel to the interface in this imaging condition. Besides the martensitic regions, regions which essentially had the structure of the parent B2 phase were also seen. Fig. 5d shows a SAED pattern
from such a region. In addition to the spots due to the B2 Phase, it was possible to see diffuse intensity distribution or streaks which were indicative of the presence of some kind of instability in the B2 phase. Fig. 5e is an HREM image of the B2 phase corresponding to the [110] zone showing the (110) planes of this phase.
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Fig. 5.
3.2.2. Specimens synthesized with a Ti particle size of ⴚ20 mm and without preheating the green compact The nature of combustion in this case was of the spin type and not oscillatory. TEM examination of the specimens showed very fine grains (Fig. 6a), SAED patterns from which indicated that these were of the ordered B2 phase. It should be noted that the extremely small size of these grains could only have been produced by rapid solidification of the alloy. No martensitic phase could be noticed in the B2 grains. The microstructures at the centre of the band and at the edge were similar. In addition to the B2 phase, it was possible to see some ellipsoidal-shaped particles in this microstructure. One typical SAED pattern from such a particle is shown in Fig. 6b. It was possible to index this pattern and many others taken from these particles in terms of the Ti4Ni2O phase.
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Continued
3.2.3. Specimens synthesized with a Ti particle size of ⴚ45 mm and after preheating the green compact Fig. 7 shows a typical optical micrograph from such a specimen. It can be seen that the microstructure comprised dendritic grains, surrounded by a eutectic phase mixture at the grain boundaries. This microstructure was somewhat similar to that observed by Zhu et al. [22] in this alloy. A detailed TEM examination of this microstructure has been carried out and the results of these investigations will be presented elsewhere [24]. The eutectic phase mixture at the grain boundary consisted of the Ti4Ni2O and the B2 phases. Some pores could also be seen in the microstructure, which is an inherent feature of the microstructure produced by micropyretic synthesis. The pores were of two types: (1) Very small pores present at the grain boundaries as well as in the grain interiors; and
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Fig. 7. Optical micrograph depicting the microstructure of combusted specimens prepared with ⫺45 µm Ti particles and preheating of compact.
Fig. 6. (a) Bright field micrograph showing fine equiaxed grains in unpreheated specimens made with a Ti particle size of ⫺20 µm. (b) [110] SAED pattern from the Ti4Ni2O phase.
(2) Much larger pores present mostly at the grain boundaries. The grain size of these specimens was much larger than that of the other three types of specimens. 3.2.4. Specimens synthesized with a ⴚ45 mm Ti particle size without preheating the green compact The microstructure of specimens obtained by this type of combustion was found to contain more or less equiaxed grains of the B2 phase. SAED patterns from these grains indicated that the B2 phase was fully ordered. A disordered bcc phase has been observed in the case of NiTi synthesized by rapid
solidification [25]. When the equiatomic NiTi alloy is slowly cooled from temperatures of the order of 1273 K, the ordering reaction occurs at temperatures between 973 K to 873 K. Since this reaction is diffusion controlled, the effect of rapid solidification will be to partially suppress it. In this case no evidence for a partial suppression of order could be found. The examination of a large number of diffraction patterns form B2 regions showed diffuse intensity maxima indicative of the presence of an omega-like instability. HREM examination of the matrix B2 phase carried out along the [110] zone yielded results which were very similar to those obtained in the case of specimens synthesized without preheat and using ⫺20 µm Ti particles. The Ti4Ni2O phase also was present in these specimens (Fig. 8a). The faceting of the particles of this phase was much less evident in this case. Fig. 8b shows a BF micrograph of a matrensitic region in this microstructure. The ⬍011⬎ type II, {111} type I, and {011} type I twins could be observed in the martensitic regions. The scale of the martensitic microstructure was finer than that observed in the case of the conventionally produced alloy. HREM of the twin interfaces in the martensitic region showed a broad strain contrast.
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[111]Ti11Ni14 / / (111)M The lattice parameter of the phase Ti11Ni14, as determined by SAED, was found to be close to 0.730 nm. Simulated diffraction patterns with this value of the lattice parameter and corresponding to the SAED patterns in Fig. 9b and c are shown in shown Fig. 9d and e, respectively. Fig. 9f shows a complex diffraction pattern from a region which had at least two variants of this phase. The key to this diffraction pattern (Fig. 9g) indicates how this pattern had resulted from the two variants. Some of the spots in Fig. 9f are due to double diffraction. Diffraction along with dark field (DF) microscopy revealed that the parallel lamellae in any given region belonged to the same variant, whereas lamellae belonging to different variants were not parallel to each other. Fig. 9h shows a typical interface between the B2 phase and the Ti11Ni14 phase. The interface was not flat and had a wavy nature. Fig. 9i shows a HREM micrograph of a region containing two variants (marked as A and B) of this phase. 3.3. Microstructural characterization of the reaction front
Fig. 8. (a) Particles of the Ti4Ni2O in the matrix. (c) Bright field micrograph showing the martensitic phase.
These interfaces were flat and did not show any ledge or unevenness. In this regard the observations made here were very similar to those made in specimens made with ⫺20 µm Ti powder after preheating the green compact. Besides the martensitic regions, there were regions which had a lamellar microstructure (Fig. 9a). SAED patterns from such a region are shown in Fig. 9b and c. These patterns could be interpreted in terms of the B2 phase and the Ti11Ni14 phase (rhombohedral unit cell with a = 103.65). The orientation relationship between this phase and the B2 matrix was found to be: (3¯ 21)Ti11Ni14 / / (1¯ 10)M
The mechanism of synthesis was examined in the case of unpreheated specimens synthesized with ⫺45 µm Ti powder. The combustion front was frozen by using a technique described in Ref. [16]. In order to examine the nature of the reaction occurring at the front, the microstructure of the specimen was examined (1) away from the front on the powder side, (2) at the front and (3) on the intermetallic side of the front. Fig. 10a shows the frozen combustion front. Depending upon the extent of reaction, the specimen could be divided in to three distinct regions: unreacted region, partially reacted region and fully reacted region. A description of the microstructure of these three regions follows. 3.3.1. Unreacted region Fig. 10b shows the powder side of the front at a distance at which the microstructure was not affected by the heat generated at the front. The Ti and the Ni powders could be readily distinguished
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Fig. 9. (a) Bright field micrograph showing a lamellar microstructure. SAED pattern from the lamellar region, (b) zone axis [111] B2, and (c) zone axis [1¯ 12] B2. (d) Simulated and superimposed SAED patterns. Open circles represent B2 reflections and filled circles represent Ti11Ni14 reflections. (e) Simulated and superimposed SAED patterns. Black circles represent B2 reflections and open circles represent Ti11Ni14 reflections. (f) Complex SAED pattern showing diffraction spots due to B2 phase, two variants of the Ti11Ni14 and spots due to double diffraction. (g) Simulated and superimposed SAED patterns of B2 and Ti11Ni14 phases. Black circles represent B2 reflections, open circles and filled circles represent reflections of two variants of Ti11Ni14. (h) HREM micrograph showing the interface between the B2 and the Ti11Ni14 phase. (i) HREM micrograph showing the B2 and two variants of the Ti11Ni14 phase (designated A and B).
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Fig. 9.
by comparing the secondary electron image with X-ray mapping. The Ti powder particles were all surrounded by Ni particles and instances of a Ni particle touching another Ni particle were very few. Ni was the lower melting ingredient and therefore was likely to melt first and surround the particles of Ti during the initial stages of synthesis. The Ti-Ni reaction is known to have high adiabatic temperature and enthalpy of formation and a low ignition temperature [10]. A study investigating the mechanism of synthesis in Ti-C-Ni-Mo alloys has shown that the Ti
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Continued
particles undergo a phase transformation wherein the a phase transforms to the b phase and then on cooling transforms to the a phase martensitically [26]. In the present study a close examination of Ti particles close to the reaction front revealed that a transformation of this nature did occur. This is clearly shown in Fig. 10c, where it could be seen that many of the Ti particles contained the martensitic phase in them. This observation was indicative of the fact that in these regions the temperature had increased beyond the a→b transition temperature (1136 K). The presence of the mar-
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Fig. 9.
tensitic phase could not be seen in particles in regions further removed from the reaction front because these had not been heated sufficiently for their temperature to have crossed the a→b transition temperature. In this description even the region in which martensitic transformation had occurred in the Ti particles has been categorized as a part of the unreacted region because the particles had not reacted with each other to any appreciable extent. 3.3.2. Partially reacted region The stopped combustion front comprised the region where the final microstructure had formed (Fig. 10a). One side of it was a region where Ni had melted and reacted with Ti either fully or partially. The melting of Ni occurred at some distance ahead of the front. This region extended from the place where the final microstructure could be seen to the point where the Ni had just melted without undergoing any substantial reaction. On moving
Continued
further towards the fully reacted region it was observed that Ni melted first and enveloped the Ti particles (Fig. 10d), leading to the formation of a Ni rich intermetallic phase which had a stoichiometry very close to Ni4Ti as revealed by EDS. An EDS profile depicting this is shown in Fig. 10e. The full scale reading in this and the other spectra shown in this work is 5000 counts. In addition, the formation of an intermetallic containing a higher amount of Ti and having a stoichiometry close to Ni2Ti could be seen (Fig. 10f). The regions from where the EDS spectra shown in Fig. 10e and f have been taken are marked as regions A and B in Fig. 10d. EDS analysis also showed that there were regions where the large Ti particles had an unreacted core (region marked as C in Fig. 10d). On moving still further towards the fully reacted region, it was possible to see two types of regions which had stoichiometries close to Ti3Ni and NiTi, respectively. These regions are marked in the micrograph shown in Fig. 10g. None of the Ti par-
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Fig. 10. (a) Micrograph showing the frozen combustion front. (b) Micrograph showing the powder side of the front at a distance at which the microstructure is not affected by the heat generated at the front. (c) Micrograph showing Ti and Ni particles. (d) Micrograph showing the melting of Ni which envelops the Ti particles and leads to the formation of a Ni rich intermetallic phase, which has a stoichiometry very close to Ni4Ti. (e) EDS profile depicting the formation of phase having stoichiometry close to Ni4Ti. (f) EDS profile showing the formation of intermetallics containing higher amounts of Ti and having stoichiometry close to Ni2Ti. The points from where EDS profiles (e) and (f) have been taken are marked as A and B in the micrograph (d). Point C in this micrograph shows an unreacted Ti core. (g) Micrograph showing an area very close to the fully reacted region. Points from where profiles showing stochiometry of Ti3Ni and close to NiTi are have been taken are marked as D and E in this micrograph.
could be construed. The molten Ni dissolved the smaller Ti particles to form Ni rich intermetallic compounds, such as Ni4Ti. The larger particles of Ti were not consumed quickly. Instead, these formed intermetallics such as Ti3Ni2 at the rim and Ti4Ni at the centre of the particle. The Ni-rich intermetallics then reacted with more of Ti and Ti rich intermetallics to form the equiatomic NiTi phase. It should be noted that in addition to the NiTi phase, the formation of the Ti2Ni or the Ti4Ni2O phases also occurred. The presence of this phase could be seen at the grain boundaries in the fully reacted side of the specimen.
ticles showed the formation of the martensitic phase. presumably because the Ti particles had undergone alloying and therefore the a→b transition temperature corresponding to these regions had changed substantially. Based on the observations made on the partially reacted region, the following sequence of microstructure formation
3.3.3. Fully reacted region The microstructure of this region was very similar to that of the fully reacted specimen in which the combustion front had not been stopped. Here a two phase microstructure could be seen (Fig. 11a). Typical EDS profiles from the center of the grain and the grain boundary, shown in Fig. 11b and c, respectively, indicated that while the composition of the center of the grain was close to NiTi, that of the grain boundaries corresponded to Ti2Ni. The size of the NiTi grains close to the front was found to be smaller than that of those away from the front.
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Fig. 10. Continued
4. Discussion 4.1. Effect of particle size and preheat on the process of synthesis It has been observed that the particle size influences the combustion process by affecting the combustion velocity [10]. The combustion velocity remains constant over a range of small particle sizes. An increase in the particle size beyond this range leads to a drastic reduction in the velocity of propagation, followed by a region where the decrease in the velocity is very weakly related to the particle size [10]. In most of the earlier studies, the Ni particle size has been altered [13]. The generation of heat by the combustion reaction and the rate of heat loss from the reaction front govern the rates of propagation, wave stability and the maximum combustion temperature during micropyretic synthesis [13]. Heat is dissipated away from the reaction front by heat transfer to the
adjacent reactant mixture, the temperature of which is below the ignition temperature, and by heat loss to the environment. A change in the reaction velocity and stability of the front will occur if these factors undergo any perturbation. Highly exothermic reactions are self-sustaining and as the enthalpy of reaction decreases, a limit is reached at which the combustion is extinguished. A nonsteady state of the reaction front results if the perturbations persist. The many kinetic reasons for a departure from steady-state propagation include particle size, compaction pressure and presence of diluents. Instability of the combustion front can arise from either inadequate heat generation during the combustion process or from a very high density of the green compact. Whereas the former leads to an instability of the front due to the non-availability of sufficient heat to sustain it, the latter causes instability because the high conductivity of the green compact leads to a rapid dissipation of heat, leading to a thermal imbalance. In the present
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Fig. 11. (a) Micrograph showing a two phase microstructure in the fully reacted region. (b) and (c) EDS profiles from the center of the grain and the grain boundary respectively.
study, the combustion process was found to be unstable when ⫺45 µm and ⫺20 µm Ti powders were used. The instability of the combustion process seemed to be due to the high density of the green compact. On preheating the specimen, however, the front could be made stable. This was indicative of the fact that the instability of the combustion front was also due to the inadequacy of the heat produced during the combustion process. The microstructure generated during the synthesis process was dependent on the Ti particle size. It will be shown later that the Ti particle size had an influence on the grain morphology and the nature of the phases forming. Microstructural studies on unpreheated specimens made from two Ti particle sizes indicated some important differences. This was because of the fact that when the particle size was ⫺45 µm the
nature of the unstable combustion corresponded to the oscillatory type propagation whereas when the particle size was reduced further to ⫺20 µm the nature of the combustion was found to change from the oscillatory to the spin combustion mode [13]. The effect of the initial temperature of the specimen on the process of combustion has been investigated by many workers [13]. The combustion temperature changes with the initial temperature of the specimen. In the case of NiAl, it has been shown by Li and Sekhar that the adiabatic temperature remains constant with the only change being an increase in the molar melted fraction of the product phase from 42 to 100% [27]. In the present study, the combustion front was found to have become stable upon preheats of the specimen and its microstructure was found to undergo a transition from equiaxed to dendritic with a substantial variation
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in the internal structure of the grains. The effect of preheat appeared to be the same as seen with other intermetallic compounds [16,17,28]. 4.2. Effect of process variables on the microstructure The microstructure was found to be considerably influenced by the two process variables, namely, preheat and Ti particle size and the effects of these variables were noticed on the morphology of the grains, the nature of the phases present and the scale of the microstructure. The important differences in the overall grain morphology and phase content are summarized in Table 2. The observation of a very fine grained microstructure in the case of specimens synthesized without preheat and with ⫺20 µm Ti powder was indicative of a high cooling rate. The grain size and the substructure were finer than those in the specimens described in Section 3.2.4, indicating that the spin type of combustion led to a higher Table 2 Effect of process variables on the overall microstructure
cooling rate than that encountered in the oscillatory mode of combustion of the alloy. The grain morphology transformed to a dendritic type on preheating of specimens synthesized with ⫺45 µm Ti. Dey and Sekhar [16] have examined the effect of preheating on the microstructure of NiAl alloys and have shown that the extent of preheat influences not only the morphology of the grains, but also the process of solute segregation. They have shown that in binary NiAl, preheating transforms the microstructure from the equiaxed to the dendritic. In NiAl alloys, containing small alloying additions such as Fe, Cr and V, it has been observed that preheating prior to synthesis results in the formation of second phase particles at the grain boundaries [16]. Such a change in microstructure and segregation of alloying elements to the grain boundaries occurs in the preheated specimens because in these the amount of the melted fraction is high and the specimens cool slowly after the combustion process, leading to solidification at a slow rate. In the case of the NiTi alloy, the pre-
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heating of the specimens had lead to the formation of dendritic grains. Moreover, segregation of the elements had occurred, facilitating formation of the eutectic microstructure at the grain boundaries. A comparison was made between this microstructure and that produced by conventional solidification of NiTi ingots and a remarkable similarity was noticed because of the fact that in either case the alloy was heated well above its melting temperature and then allowed to cool slowly. In specimens with a Ti particle size of ⫺20 µm, the grains were equiaxed even after preheating. This was so because heat dissipation occurred at a more rapid rate as compared to the case of the specimens with a powder size of ⫺45 µm. It was interesting to note that the B2 phase showed almost perfect order even in specimens, which had been synthesized with an unstable front. Though this phase formed by a rapid processing technique involving the reaction of the elemental powders of Ti and Ni, the mixing of the elements was quite homogeneous. The SAED patterns from the B2 phase showed streaks (Fig. 5d). Such streaks have also been noticed in the B2 phase of the NiTi alloy produced by the conventional route [2] and are quite similar to those observed in diffraction patterns obtained from b phase of the w forming systems [29]. The b→w transformation has been mostly found to occur in the group IV metals [30]. Experimental work as well as theoretical studies have been carried out to explain the diffuse w scattering [30]. HREM studies have been performed by Schryvers and Tanner [29] to ascertain the changes occurring in the b lattice when the diffuse w scattering initiates. Their HREM studies have shown that the diffuse w scattering is due to the presence of short strings of one and two-dimensional w phase. HREM was carried out in the present study along the [110] B2 zone axis. It could be seen that the HREM image (Fig. 5e) did not show the presence of features like the ω strings observed by Schryvers and Tanner [29]. The diffuse maxima seen in the diffraction pattern, therefore, could not be ascribed to the type of structural transformation observed in the case of w forming alloys. In the case of NiTi, the formation of an ω type phase after aging of the B2 phase has not been reported. A detailed HREM study of the diffuse
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intensity distribution in the diffraction patterns from the B2 phase in NiTi has not been carried out even in a conventionally produced NiTi alloy. Moine et al. [31,32] have examined the premartensitic diffraction effects which occur just above the martensitic transition temperature by examining the electron diffraction patterns. They have attributed the observed extra reflections (not diffuse intensity maxima) to the occurrence of two types of particles and the presence of lattice displacement waves in these particles. In this study, however, no extra reflections in the SAED patterns from the B2 phase could be noticed. The structure of the martensitic phase observed in this study was very similar to that produced by conventional processing. This could be ascertained both by conventional microscopy, as well as by HREM. HREM examination of the twin interfaces indicated that these were flat and did not show any ledges or unevenness. The observations made in this study on the structure of the interface were similar to those made by Nishida et al. [5] and were at variance with those made by Knowles [1]. The scale of the microstructure, however, was much finer than that in the case of conventional processing. The Ms temperature of the specimens examined in this study was found to be higher than that of the conventionally produced NiTi alloy. It has been mentioned earlier that the microstructure produced under conditions of unstable combustion were found to have features, such as very fine grain size, which suggested the occurrence of solidification under a high cooling rate. Therefore it is worth comparing the Ms of specimens of NiTi synthesized by unstable combustion with those produced by rapid solidification. Igaharo and Wood [25] have shown that the Ms temperature of NiTi is lowered by rapid solidification. Factors which could affect the transformation behavior of this alloy following rapid solidification include the high concentration of quenched-in vacancies, the dislocation density, the degree of disorder and the grain size. Since the equilibrium concentration of vacancies increases with temperature in an exponential manner, it is expected that a rapidly solidified alloy will have a higher concentration of vacancies as compared with an alloy quenched from the solid
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state [25]. The vacancies formed in excess tend to form vacancy clusters, which may collapse to yield dislocation loops. It has also been seen that some of these vacancies anneal out if the cooling rate during the post solidification stage is not high enough [25]. The strain field associated with the dislocations and dislocation loops tend to promote martensite formation and consequently raise the Ms by providing a favorable interaction with the strain fields of the martensite nuclei [25]. The other factors affecting the Ms value are the grain size, the extent of order in the parent phase and the amount of oxygen present in the material. In this study it was observed that the Ms temperature for the alloy synthesized under an unstable condition was higher than that for an alloy produced by conventional solidification. This was indicative of the fact that the aforementioned factors had indeed played an important role in relation to the Ms temperature for specimens examined in this study. It has been shown by Moore and Yi [33] that a higher cooling rate from the molten state leads to a dendritic morphology of the Ti4Ni2O phase (designated in their study as the Ti2Ni phase), whereas a slow cooling rate leads to a faceted morphology of this phase. In the present study it was seen that the morphology of the Ti4Ni2O phase indeed depended on the processing condition. Besides the eutectic morphology a faceted morphology was also observed in specimens produced under a different set of processing conditions. The Ti4Ni2O phase has a fcc structure of the Fe3W3C type with a lattice parameter a = 1.13193 nm, while the Ti2Ni phase also has an identical structure with a lattice parameter a = 1.13279 nm [34]. Ti2Ni can transform to a quasicrystalline structure on rapid solidification [35]. The quaiscrystalline phase forming on rapid solidification in a (Ti1-x Vx)2Ni alloy, specially after heat treatments at 673 K, shows a kind of local or short-range transnational order as evidenced by a more or less commensurate arrangement of electron diffraction spots along one or several two-fold directions [35]. Though the cooling rate during unstable synthesis was high, it was not high enough to lead to the formation of a quasicrystalline structure in this study. The Ti11Ni14 phase could be seen in one of the specimens in an unusual plate shaped morphology.
It may be noted that this phase has been found to form in the NiTi alloy after ageing at temperatures below 973 K and has been seen to adopt an oval plate morphology with lenticular cross-section [34]. Although the crystallographic orientation relationship between this phase and the B2 matrix observed in this study was the same as that seen earlier [34] in other studies [34], the lattice parameter of the unit cell was found to be larger (a = 0.720 nm) as compared to a = 0.661 nm reported in Ref. [34]. The formation of this phase in a lamellar arrangement, noticed in the present study has not been reported so far. This phase plays an important role in the development of the all round shape memory effect in the NiTi alloy [2]. Even though the morphology observed in this study resembled that of a martensitic phase, the interface structure and the substructure showing neither a high density of dislocations or of twins appeared to discount this possibility. The features of the Ti11Ni14 precipitates were suggestive of their having been formed by a diffusional transformation. This phase could be seen only in specimens synthesized with ⫺45 µm Ti. 4.3. Mechanism of synthesis The examination of the mechanism of synthesis could help in establishing the sequence of events leading to the formation of the final microstructure and in identifying those process parameters which control the microstructure. In this study the mechanism of synthesis was studied by stopping the combustion front, by making use of an arrangement of copper plates around the specimen, as described in Ref. [16]. Microstructural examination was carried out at different regions of the specimen after stopping the combustion front. Based on the microstructural examination of the three regions in the partially combusted specimen, the following sequence for the formation of the final microstructure from the powder particles could be proposed. The formation of the NiTi phase is triggered by the melting of the Ni particles, which then engulf the Ti particles. At first, Ni rich phases such as Ni4Ti and metastable phases form. The Ni rich phases then react with more of Ti and yield phases which are progressively richer
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in Ti. Ultimately, as the temperature increases, Ni reacts with Ti2Ni3 and Ti to form stoichiometric TiNi. Some amount of the Ti2Ni phase remains unreacted and becomes a part of the final microstructure. The mechanism of synthesis described here is pertinent to one set of processing conditions (Ti particle size and preheat). The mechanism may differ in minute details when examined for another set of processing conditions because the final microstructure depends on these conditions. However, the basic steps in the mechanism could be expected to remain the same.
5. Conclusions On the basis of the observations made in this study, the following conclusions could be drawn: 1. It was possible to generate a variety of microstructures in NiTi synthesized by micropyretic synthesis by appropriately choosing the processing conditions. These microstructures could range from the dendritic microstructures commonly encountered in ingots in the as-cast condition to equiaxed microstructures obtained after thermomechanical processing. Very fine grained microstructures could be obtained in regions undergoing unstable combustion. 2. The combustion process was unstable when preheating was not used. This instability appeared to be due to the fact that very fine Ti and Ni powders were used and the thermal conductivity of the resulting compacts was very good, promoting rapid heat dissipation. The nature of the instability was found to change from the oscillatory mode of combustion to spin type combustion with a decrease in the Ti particle size. This was presumably due to a progressive increase in the thermal conductivity of the compact with decreasing Ti particle size. 3. HREM examination of the B2 phase indicated that the instability in this phase, manifested in the form of diffuse maxima in SAED patterns, was not due to the formation of an w-like structure. 4. The nature of the martensitic phase was the same as that encountered in the conventionally
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processed alloy except for the fact that the size of the martensite plates was much finer. The martensite twin boundaries were devoid of any ledges or steps and were very similar to those observed in the conventionally produced alloy. 5. The Ti4Ni2O phase had the same structure and morphology as observed in the case of the conventionally produced alloy. The morphology of the phase was found to depend on the processing conditions. The Ti11Ni14 phase was found to show an unusual lamellar morphology. 6. The mechanism of synthesis of the alloy involved melting of Ni, followed by reaction with Ti to form Ni rich intermetallic compounds. These then reacted with more of Ni and Ti to give rise to intermetallic compounds with higher and higher amounts of Ti, finally leading to the formation of the NiTi phase.
Acknowledgements The author would like to thank Dr. K. Madangopal, Dr. D. Srivastava, Dr. P. Mukhopadhyay and Dr. P. K. De for many helpful discussions. Thanks are also due to Prof. J. A. Sekhar for his constant encouragement during the course of this work.
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