Martensitic Transformation

In addition, Yoshikawa et al. [12] investigated the martensitic transformation by 0.02% proof stress of Fe- SMA under various loading condition. As the result, the martensitic transformation behaviour cannot be distinguished under various low stress loading condition. However, quantification of stress-induced martensite has not been done. In the past, Stanford and Dunne [4] conducted the bending tests of Fe-SMA. As a result, it reported that the region subjected to compressive stress has more rumpled surface relief than the region subjected to tensile stress. It proves the the change in volume resistivity under compression becomes larger than that of tension as shown in Fig. 11.

by diffusion process, both aging temperature and time will greatly affect the transformation behavior. Therefore, it is very interesting and useful to find out how transformation proceeds with different aging treatments. The purpose of the present study is to investigate the effect of aging on succes- sive martensitic transformation in near-equiatomic Ti-rich Ti– Pd alloys. Subsequently, the equilibrium phase boundary of TiPd compound in Ti-rich side is estimated on the basis of transformation behavior.

Fe–24 at%Pt and Fe–29.9 at%Ni alloys were produced by melting the component metals in a high frequency induc- tion furnace under argon atmosphere and casting into a wa- ter cooled iron mold. An ordering heat treatment was made in the Fe–24 at%Pt and its degree of order was 0.8 (the or- dered Fe–24 at%Pt alloy thus obtained exhibits a thermoelas- tic martensitic transformation at M s = 173 K). Single crys-

In this study, an identical area of a SUS304 austenitic stainless steel specimen was observed by electron backscattering diffraction measurements at different strains in tensile test at ambient temperature, in order to investigate the details of deformation-induced martensitic transformation. Firstly, a number of thin deformation twins were formed in austenite grains. Most of the martensite crystals were observed either near grain boundary triple junctions or inside the deformation twins. Secondly, it was found that martensite crystals preferentially appeared in the austenite grains whose h001i crystal directions were close to the tensile direction. Furthermore, when austenite grains had several martensite crystals inside, only one or two variants were observed among 24 variants possible under Kurdjumov-Sachs orientation relationship, which indicated the existence of variant selection rules. Patel and Cohen model and Bogers ­ Burgers model were examined to understand the variant selection, but both models could not explain the variant selections. The result suggests that complicated stress states govern the deformation- induced martensitic transformation in polycrystalline austenite. [doi:10.2320/matertrans.MBW201212]

On the other hand, the obtained martensite in the specimen slowly cooled to 4 K with a cooling rate of 0.17 K/s was only 7% in fraction as shown in Fig. 1. Apparently, the trans- formability by simple cooling exhibits cooling-rate depend- ence. Furthermore, this result means that the martensitic transformation of Fe-Co particles is prohibited if the cooling rate is sufficiently small to prevent the introduction of thermal stresses. We now confirm that the following tests in the present study enable us to reveal the net effects of the applied magnetic field on the martensitic transformation of the Fe-Co particles.

Abstract: Zr-Cu binary and Zr-Cu-Al ternary alloys with different compositions were fabricated using arc melting. The phase structure and martensitic transformation temperatures of the alloys were investigated using X-ray diffraction and differential scanning calorimetry, respectively. It was found that the ZrCu martensitic phase was formed as an intermetallic compound in the near-equiatomic Zr-Cu binary alloy. On the other hand, both the ZrCu martensitic and parent phases were formed in the Zr-Cu-Al ternary alloy. In addition, it was confirmed that the martensitic transformation temperature of ZrCu decreased with addition of Al to the base alloy. Therefore, it was found that addition of Al to the equiatomic Zr-Cu alloy can effectively control the microstructure and martensitic transformation temperature.

seen that the monolithic FeNi film exhibits a fcc struc- ture (γ), without existence of martensite (α) with a bcc structure. With the initial increase of V layer thickness, bcc-structured FeNi phase is detected in nanomulti- layered films besides fcc-structured FeNi phase, suggest- ing that martensitic transformation occurs in FeNi layers. In addition, with the increase of deposited time from 2 to 6 s, the diffraction peaks for fcc-structured FeNi weaken, while those for bcc-structured FeNi strengthen. According to the deposition rate of V (about 0.25 nm/s) derived from the monolithic V film, the thicknesses of the V layers deposited for 2, 4, 6, 8, 10, and 12 s at the same condition are 0.5, 1.0, 1.5, 2.0, 2.5, and 3.0 nm, respectively, which have been indexed in the corresponding XRD patterns in Figure 2. When the V layer thickness increases from 1.5 to 2.0 nm, however, the bcc-structured FeNi can hardly be detected, implying that the martensitic transformation of FeNi terminates. As the V layer thickness further rises to 3.0 nm, the (110) diffraction peak of bcc-structured V emerges in the XRD patterns besides fcc-structured FeNi, suggesting that V layers begin to present a stable bcc structure.

1.6) alloys is investigated. The addition of Gd results in a change in the microstructure. With the increase of Gd content, the grain size is clearly reduced and the volume fraction of the second phase increases gradually. And then the second phase along the grain boundaries grows and connects to each other. Some disperse in the grains. One-step thermoelastic martensitic transformation is observed in Ni 45 Co 5 Mn 35 In 15−x Gd x alloys. When the Gd content reaches 0.8 (at%), the martensite transformation temperature decreases first, and then

determination of constitutional phases was carried out by using an X-ray diffractometer (XRD). DSC measurements were performed with a cooling and heating rate of 10 K/min in order to estimate the martensitic transformation temper- atures. The temperature range measured was about 130 to 800 K. The TEM specimens were electropolished using a twin jet method in an electrolyte solution consisting of 25% HNO 3 and 75% CH 3 OH by volume at around 238 K. TEM

cooling process. In the heating process, the –T curve starts to increase at about 60 K, meaning that the 0 -martensite was also formed in the heating process. This result implies that the martensitic transformation of the sensitized SUS304 proceeds isothermally in the temperature range between 60 and 260 K. Furthermore, we notice that the –T curve increase in two step in the cooling process as indicated by ‘‘A’’ and ‘‘B’’, respectively. This result suggests that there are two kinds of isothermal martensitic transformations in the sensitized SUS304 stainless steel.

3.1 Effect of temperature on transformation behavior In order to investigate the martensitic transformation behavior and magnetic properties, we have measured temper- ature dependence of magnetic susceptibility by applying a low magnetic field of 79.4 kA/m. Figure 1(a) shows -T curves of the SUS304 in the solution-treated and sensitized states. The -T curve of the solution-treated SUS304 shows a sharp peak at about 40 K due to a paramagnetic to antiferromagnetic transition of the -phase, being in agree- ment with a report by U. Gonser et al. 13) There is no

decisive choice of parameters because there is still contro- versy about exact crystalline structure of Ni–Ti M-phases, as described above. There is also an essential reason that in actual materials there must be a lot of microscopic combi- nations of local crystalline configuration depending on chemical stoichiometry (e.g. useful SMA is not Ni at% = 50 and Ti at% = 50). However, we could consider the atomistic mechanism of Ni–Ti by using possibly precise potential such as embedded atom method (EAM) type, if we have confirmed applicability of potential with some evalua- tion. Indeed, there is a success of adopting Lennard–Jones type (pairwise type) potential for Ni–Ti to the study of martensitic transformation mechanism. 4) From the point of view of energetics, utilizing EAM is supposed to be better. Modified EAM (MEAM) should be recognized to be a future possibility for interactions in SMA systems, though pursued with extensive effort by other researchers. 12)

for static loadings. In order to follow this phenomenon at a tension velocity of 1 m/s, complementary measurements are conducted. Displacement and subsequent axial strain path along the specimen are calculated by digital image correlation (DIC). Moreover, since martensitic transformation is exother- mic, infrared thermography (IRT) measurements are also a good mean to measure indirectly the loca- tion of the ongoing transformation. Optical and thermal measurements are obtained simultaneously on the same thin specimen, on two opposite sides that are supposed to have the same thermo-mechanical response. These two optical techniques are described in following sections, along with the experimen- tal strategy that has been implemented.

from 400 to 103 K. To investigate precisely the process of the martensitic transformation, changes in Laue spots during the transformation process were observed. The following phenomena were discovered: The single crystal at 300 K before increasing temperature shows a tetragonal structure, the a- or b-axis of which is parallel to the direction of crystal growth. It shows a cubic Heusler structure at 400 K after heating. One of the cubic axes is parallel to the direction of crystal growth. With decreasing temperature, the crystal transforms into many small variants with tetragonal structures. One of the variants evolves so that the a- or b-axis is parallel to the direction of the crystal growth, and the orientation of the other variants can be represented by its rotation about the cubic h111i axis. The rotation angle is at most a few degrees, which changes gradually with decreasing temperature. At 103 K the transformation is almost complete and many tetragonal variants are rearranged into one tetragonal structure, the a- or b-axis of which is parallel to the direction of crystal growth.

Recently, the present authors have reported that ductile Co-Al alloys containing Al over 10 at% exhibit an SM effect associated with (fcc)/" (hcp) martensitic transformation and that the reverse transformation temperatures are located between 200 and 300 C. 15) One of the characteristic features

Figure 5 shows the dependence of the 0.2% proof stress of the three investigated alloys on deformation temperature. The Fe-33Mn alloy in which ¾-martensite does not form showed a monotonic decrease in 0.2% proof stress with increasing deformation temperature. This reduction in 0.2% proof stress is explained by the promotion of slip deformations associated with thermally activated dislocation motions. In contrast, the Fe-33Mn-4Si and Fe-33Mn-6Si alloys showed increasing 0.2 % proof stress with increasing deformation temperature in the range from 273 to 373 K, but decreasing values with further increasing temperature in the range from 373 to 423 K. Although the latter trend is also explained by the thermal activation process of dislocation motions, the former trend requires the consideration of the critical stress for the deformation-induced ¾-martensitic transformation. The crit- ical stress for the deformation-induced ¾-martensitic trans- formation is controlled by the difference in enthalpy between the £ and ¾ phase. According to the Clausius-Clapeyron equation, the 0.2% proof stress increases with increasing deformation temperature as long as yielding is dominated by the initiation of the ¾-martensitic transformation (for a more quantitative understanding, this consideration needs to be modified according to a recent work 32) ). In other words, the

EDS, hardness and cold-rolling workability tests. Experimental results show that all these HTSMAs exhibit B2 ! B19 martensitic transformation with Ms temperature can be predicted by the linear regression as: Ms ðKÞ ¼ 424:2 þ 2:7X 9:7Y. The substitution of Ni/Pd by Cu affects the lattice constants of B19 martensite, the hardness and the cold-rolling workability which can be explained by the different atomic radii of Ni, Cu and Pd. Ti 50 Ni 15 Pd 25 Cu 10 HTSMA annealed at 523 K up to 9 weeks and thermal-cycled up to 100 times still has quite good

In the present study we only investigate the aging of fully annealed, defect free materials. The multistage martensitic transformation has been also observed in the aged specimen after deformation. 8,10,11) Since it is likely that the solution treatment atmosphere prior to deformation was not carefully regulated in the previous works, 8,10,11) it should be reinvesti- gated with respect to the heat treatment atmosphere. How- ever, it is difficult to regulate the heat treatment atmosphere from the industrial point of view. Therefore, further under- standing of the multistage martensitic transformation mech- anism is also required, although the transformation is an extrinsic nature, i.e., a kind of artifact during the heat treatment. Experimentally unsolved questions derived from the present study are listed as follows. The compositional fluctuation between grain boundary and interior upon solution treatment under normal atmosphere such as the condition A-A should be estimated with aid of analytical TEM, Auger electron spectroscopy and so on. Then the heterogeneous nucleation and growth mechanism of Ti 3 Ni 4

and shape memory effect of rapidly solidified TiNi alloys with some ternary elements have been investigated by some researchers. 17,18) A common conclusion of these studies is that microstructures, martensitic transformation behavior, mechanical properties and shape memory properties of rapidly solidified TiNi alloys are quite different from those of bulk material. For examples, rapidly solidified Ti 50 Ni 25 -

occurrence of martensitic transformation. Figure 1 impli- cates that the number of dislocations induced in the alloy is getting higher with the increasing number of thermal cycling in the first 20 cycles. The reason why two transformation peaks can be induced by thermal-cycling, as shown in Fig. 1, will be discussed in the following.