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Phase equilibria in the Ni–Mn–In alloy system
Journal of Alloys and Compounds 549 (2013) 57–63
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Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Phase equilibria in the Ni–Mn–In alloy system T. Miyamoto, M. Nagasako ⇑, R. Kainuma Department of Materials Science, Graduate School of Engineering, Tohoku University, 6-6-02 Aoba, Sendai 980-8579, Japan
a r t i c l e
i n f o
Article history: Received 12 July 2012 Received in revised form 29 August 2012 Accepted 30 August 2012 Available online 6 September 2012 Keywords: Phase equilibria Diffusion triple Ni–Mn–In Martensitic transformation Order–disorder transition Magnetic transition
a b s t r a c t Phase equilibria at 700 and 850 °C, critical temperatures of B2/L21 order–disorder transformation and martensitic and ferromagnetic phase regions at room temperature in the Ni–Mn–In system were determined mainly by diffusion triple method using a two-stage diffusion couple technique. It was confirmed that a single phase region of the b phase at both 700 and 850 °C exists in a wide composition range along the NiMn–NiIn section and that the L21 ordered phase region appears in the vicinity of Ni2MnIn in the temperature region below about 800 °C. The composition lines, iso-Ms and iso-TFM, possessing the Ms and TC temperatures at room temperature, respectively, were successfully estimated and the coincidence between the iso-Ms and iso-TFM was confirmed in the composition region from 10 to 20 at.% In in the bphase region. Ó 2012 Elsevier B.V. All rights reserved.
1. Introduction Since unusual martensitic transformation from the ferromagnetic austenite phase to the paramagnetic martensite phase has been reported in the off-stoichiometric Ni2MnIn and Ni2MnSn alloys with Heusler L21-type ordered bcc structure [1–4], Ni–Mn–In based alloys have received much attention as high-performance multifunctional materials. The martensitic transformation temperatures of these alloys are drastically decreased by application of a magnetic field and magnetic-field-induced reverse transformation, namely, metamagnetic phase transition, occurs near the martensitic reverse transformation start temperature, Ms [2,3]. Moreover, a magnetic-field-induced shape memory (SM) effect, i.e., metamagnetic SM effect, was confirmed in Ni45Co5Mn36.7In13.3 [4] and Ni43Co7Mn39Sn11 [5] alloys at room temperature, and many interesting phenomena accompanying this unique transformation, such as the inverse magnetocaloric effect [2,6], the giant magnetoresistance effect [7,8] and giant magnetothermal conductivity [9] were reported. While the basic physical properties of the Ni–Mn–In Heusler alloys have been reviewed by Planes et al. [10], no phase diagram covering the whole range of the Ni–Mn–In ternary system has been established. Fig. 1 shows the most reliable Ni–Mn [11,12], Mn–In [13] and Ni–In [14] binary phase diagrams to date. It is seen that the single phase region of the a (A1: disordered fcc Ni or Mn) exists in a wide composition range in the Ni–Mn binary system and that the b (B2: CsCl-type ordered bcc NiMn or NiIn) phase in relation to the
⇑ Corresponding author. Tel./fax: +81 22 795 5256. E-mail address: [email protected] (M. Nagasako). 0925-8388/$ - see front matter Ó 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2012.08.128
Ni2MnIn Heusler alloy appears in the temperature range from 700 to 1000 °C in the vicinity of the stoichiometric composition in the Ni–Mn and Ni–In systems. From this information, the b-NiMn phase region is expected to be connected to the b-NiIn phase region at high temperatures. In spite of the importance of further alloy development, only limited information on the phase equilibria at elevated temperatures in this ternary system has been reported. In the present study, the phase equilibria at 700 and 850 °C in whole region of the Ni–Mn–In ternary system are determined and the composition region where martensite and ferromagnetic phases are observed at room temperature is evaluated by using the diffusion triple (DT) method. Here, the DT method is a kind of combinatorial methods, first reported by Jin [15] and more recently employed by Zhao et al. [16]. 2. Experimental procedures 2.1. Diffusion triples Details of the DT method that we used in the present study were reported in our previous paper on the Co–Fe–Ga phase diagram [17]. The DT specimens for the Ni– Mn–In system were prepared by the following three steps: (1) Pure Ni (99.9%) and Mn (99.9%) blocks were joined with induction melting in an argon atmosphere, and only the Mn block laid on a Ni block was melted in the induction furnace. Ni/Mn diffusion couple (DC) samples were diffusion-annealed at 1000 °C for 168 h in evacuated quartz tubes to obtain a continuous diffusion zone between the Ni and Mn blocks. (2) A cylindrical hole with a diameter of 3 mm was formed by electro-spark machining near the diffusion zone of the Ni/Mn DC, and some In (99.99%) chips were inserted into the hole after removal of the surface layer by polishing. Diffusion triples were finally obtained by diffusion-annealing at 700 °C for 24 h, 96 or 168 h and at 850 °C for 3, 12, or 24 h in evacuated quartz tubes.
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T. Miyamoto et al. / Journal of Alloys and Compounds 549 (2013) 57–63
Fig. 1. Binary phase diagrams constituting the Ni–Mn–In ternary system [11–14].
Table 1 Equilibrium concentration in the Ni–Mn–In system determined by the diffusion triple method. Temp. (°C)
Time (hours)
Phase U
Phase C
Equilibrium composition (at.%) Phase U
850
3
f
c b 12
c c c c b
c 24
b-Mn
c b-Mn 700
24
c Ni2In3
c Mn3Ni2In b
c
96 168
b Mn3Ni2In b b b
c c f
c
Ni3In b L b f f b f b-Mn L b-Mn L b-Mn L b b-Mn Mn3Ni2In b-Mn Mn3Ni2In L Mn3Ni2In Ni2In3 Ni2In3 b f Ni3In Ni3In
Phase C
Mn
In
Mn
In
0.4 22.1 2.9 41.5 15.0 10.9 28.0 12.5 4.0 68.8 67.2 64.6 72.1 1.6 67.3 56.7 32.0 69.9 42.2 48.0 51.4 9.3 7.7 20.8 13.2 1.4 9.4
30.9 7.6 57.2 0.3 6.2 6.6 2.8 32.2 34.2 18.0 2.4 11.0 0.8 63.3 1.2 13.3 28.8 1.3 15.6 21.3 7.8 46.3 48.5 3.3 4.2 30.4 2.8
0.3 20.5 6.6 45.6 5.7 3.5 25.9 4.8 10.8 56.3 65.5 54.1 71.9 0.8 55.8 69.1 48.3 69.2 49.3 36.6 53.4 3.0 3.6 19.2 3.2 1.4 1.8
24.6 25.4 62.7 4.1 30.6 30.3 21.8 32.6 33.8 27.9 6.0 20.4 4.4 85.5 6.1 5.1 21.1 5.0 19.7 40.4 15.1 55.3 55.0 24.9 30.5 24.6 24.2
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T. Miyamoto et al. / Journal of Alloys and Compounds 549 (2013) 57–63 Table 2 Equilibrium concentration in the Ni–Mn–In system determined by the conventional alloy method. Composition (at.%)
Temp. (°C)
Time (hours)
Phase U
Phase C
Phase H
Equilibrium composition (at.%) Phase U
Mn
In
78 59 55
17 6 40 30 5 10 32 16 28 38
850
17 6 40 30 20 25 18 5 32 16 15 50 51 32 28 45 38
700
41 36 30 8 4 78 59 55
50 42 41 30 25 15 12 8 5 4
24
b-Mn
c Mn3In b-Mn
c c b
c c b 72
b-Mn
c Mn3In Mn3In Mn3Ni2In Mn3Ni2In b
c b
c c b b b
c b b
Phase C
Phase H Mn
In
Mn3In b L L b b L b b f
– L – – – – – – f –
79.8 65.7 70.2 69.8 38.6 34.5 29.0 31.1 17 9.3
15.3 2.4 26.0 16.7 0.9 1.7 26.4 2.5 6.3 37.9
70.6 55.9 47.6 54.2 41.1 35.4 26.8 29.9 6.1 3.1
25.1 8.0 45.3 29.7 6.6 11.6 43.4 17.4 30.3 35.9
– 56.3 – – – – – – 13.6 –
– 17.7 – – – – – – 28.1 –
Mn3In b L L b-Mn Mn3In Mn3Ni2In b Mn3Ni2In b b L L f b NiIn f
– Mn3Ni2In – – – L – b’ L – – – – – f – Ni13In9
79.1 67.9 65.6 58.2 52.0 51.0 40.7 42.3 29.6 31.8 28.5 16.8 13.4 15.5 17.6 10.0 11.7
15.8 1.4 28.2 26.5 20.4 22.4 16.9 1.1 30.9 1.1 1.6 41.8 45.0 29.1 4.3 40.0 35.9
73.6 53.0 30.2 36.2 64.9 57.9 49.2 42.6 45.0 29.2 25.8 27.7 6.9 5.4 17.3 0.3 2.7
21.7 7.7 62.8 42.6 16.4 25.7 18.5 6.6 23.5 19.5 22.0 50.3 73.8 31.5 25.8 48.9 35.5
– 54.3 – – – 35.9 – 39.4 32.2 – – – – – 5.0 – 1.7
– 14.9 – – – 42.0 – 0.6 42.8 – – – – – 30.3 – 39.3
(3) The obtained DTs were cut from a section vertical to the In cylinder and subsequent to polishing were chemically analyzed using an electron microprobe analyzer (EPMA) along some lines parallel to the Ni/Mn interface. Consequently, information on phase equilibria over a wide composition range was extracted from the line analyses along the diffusion paths, as listed in Table 1.
Mn
In
Mn
In
(a)
In the present study, the obtained DTs covered most of the composition range in the ternary system except for the In-rich corner in which a liquid phase appeared. 2.2. Alloy specimens The alloy samples shown in Table 2 were additionally prepared and the microstructures were analyzed by EPMA in order to confirm reliability of the phase equilibrium determined by the DTs. Ni30Mn50In20 alloy was also prepared to determine the crystal structure of a new phase detected in the DT examination. These alloys were induction-melted from pure Ni, Mn, and In, and annealed for equilibration at 700 °C for 72 h and at 850 °C for 24 h in quartz tubes under an argon atmosphere. Microstructure observation was carried out by optical microscopy (OM). The crystal structure was determined by transmission electron microscopic (TEM) examination for thin foils and by X-ray powder diffraction (XRD), where the TEM sample was prepared by jet-electropolishing with a solution of 8 vol.% perchloric acid, 72 vol.% acetic acid, 8 vol.% ethylene glycol and 12 vol.% ethanol at room temperature. The order–disorder transformation temperature was determined by differential scanning calorimetry (DSC) at a heating rate of 10 °C/min.
(b)
2.3. Determination of iso-Ms and -TFM lines Regions of alloy composition possessing martensite and ferromagnetic phases at room temperature were determined from the microstructure in the DTs. In the case of the martensite phase, since the chemical compositions of the alloys possessing the martensitic transformation starting temperature T Ms at room temperature correspond to those at the edge of the martensite region in the DTs, the iso-Ms (= RT) composition line can be estimated by chemical analysis at many positions on the edge of the martensite region, as reported in our previous papers [17]. It is impossible to distinguish between ferromagnetic and non-magnetic regions by conventional OM observation. In the present study, the ferromagnetic region was determined by contrast of magnetic walls visualized by magnetic colloids.
Fig. 2. (a) Microstructure and (b) concentration–penetration profiles for Ni and Mn elements in the vicinity of the diffusion zone in the diffusion triple annealed at 850 °C. The martensite region partially formed in the b (B2) phase region during cooling from annealing temperature is labeled with ‘m’. Equilibrium concentrations were evaluated by extrapolation of the profiles to the phase interface.
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T. Miyamoto et al. / Journal of Alloys and Compounds 549 (2013) 57–63
The chemical compositions of the alloys possessing the ferromagnetic to nonmagnetic transition temperature at room temperature, i.e., the iso-TFM (= RT) compositions, were evaluated by chemical analysis at many positions on the edge of the ferromagnetic phase region as well as the iso-Ms (= RT) compositions. Here, the edge of the ferromagnetic phase region was determined by using image analysis.
(a)
3. Results and discussion 3.1. Isothermal section diagrams of 700 and 850 °C Fig. 2(a) shows the OM image in the vicinity of diffusion zone in the DT specimen annealed at 850 °C, where three distinct regions, i.e., Ni-rich c, Mn-rich c and In-rich b phases, are observed on the left, lower right and center sides. Besides these phases, a martensite phase region with a lamella-like contrast and a liquid phase region with a solidification structure are observed as labeled with ‘m’ and ‘Liquid’ in Fig. 2(a), respectively, where the martensite phase region is formed in the b phase region and there is a phase boundary with the austenite phase region with no lamella-like contrast. Since the martensitic transformation occurs during cooling, the phase boundary dividing the austenite from the martensite region corresponds to a habit plane and has no relation to the phase diagram at the annealing temperature. Equilibrium compositions among those observed phases were determined using concentration–penetration profiles obtained by EPMA for the DTs annealed at 700 and 850 °C. Fig. 2(b) shows typical concentration–penetration profiles for Mn and Ni elements in the vicinity of the diffusion zone in the DT specimen annealed at 850 °C. Phase equilibrium concentrations were determined by extrapolation of the concentration profiles to the phase interface as demonstrated in Fig. 2(b). All the results obtained from the DT samples annealed at 700 and 850 °C are listed in Table 1. Because it is difficult to determine three-phase and solid + liquid two-phase equilibria using the DT method, some alloy samples listed in Table 2 were also prepared. A typical microstructure obtained from the Ni35Mn59In6 alloy is shown in Fig. 3. In this specimen annealed at 850 °C followed by quenching, three phases, the c, the b and the Liquid (L) phases, are observed, whereas the b phase has transformed to the martensite phase with the lamellalike contrast during cooling from the annealing temperature. Average phase equilibrium concentrations were determined from EPMA data obtained from at least seven different points for each phase. All the results obtained from the DT samples and alloy specimens annealed at 850 and 700 °C are plotted in Fig. 4(a) and (b), respectively. It is clear that phase equilibria determined by the DT method are coincident with those by the alloy method. In the isothermal diagrams of 700 and 850 °C, the topological constitutions on the phase equilibria are not complicated and are similar to each other, except for the presence of Mn3Ni2In compound at 700 °C, where the Mn3Ni2In phase is a ternary intermetallic com-
(b)
Fig. 4. Phase equilibria at (a) 850 °C and (b) 700 °C of Ni–Mn–In ternary system determined with diffusion triples and alloy specimens. Symbol (+) plotted in b phase of 700 °C diagram indicates points on critical boundary for the chemical 1 ordering estimated from the T B2=L2 curves of Fig. 6. t
pound that is firstly found in the present study. The liquid phase widely extends to the middle region of the ternary diagrams and the NiMn phase with B2 structure exists along the NiMn–NiIn section including the Heusler composition Ni2MnIn. According to the Ni–In binary phase diagram shown Fig. 1 [14], three binary compounds Ni13In9 (B81: NiAs-type), NiIn (B35: CoSn-type), and Ni2In3 (D513: Al3Ni2-type) exist up to 876, 860 and 870 °C, respectively. All these phases, however, do not appear in the 850 °C diagram. This means that the phase equilibria in the Ni–In binary system should be more accurately reexamined. On the other hand, all these binary phases appear in the 700 °C diagram, while their existing regions are limited on the binary side. The basic crystal structure of the Mn3Ni2In compound was determined by TEM observation. From analysis of obtained electron diffraction patterns, the crystal structure of this phase was determined to belong to the cubic system with space group Fig. 5(a) shows the XRD pattern taken from the Ni30Mn50Fd 3m. In20 single-phase alloy. Based on the space group and the experimental XRD pattern, the prototype of this phase is conformed as being the Mn3Ni2Si structure in which Mn, Ni, and Si (In) atoms occupy 48f, 32e, 16d sites, respectively, as shown in Fig. 5(b). The lattice constant of this phase in the Ni30Mn50In20 alloy is estimated as a = 1.1307 nm. 3.2. Order–disorder transformation in the Ni(Mn, In) phase
Fig. 3. Three-phase microstructure taken from Ni35Mn59In6 alloy annealed at 850 °C for 24 h, where the b (B2) phase has transformed to the martensite phase during quenching.
Order–disorder transformation from the B2 to L21 phase in Ni50Mn50 xInx (hereafter 50%Ni) alloys has recently been investigated
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T. Miyamoto et al. / Journal of Alloys and Compounds 549 (2013) 57–63
(a)
20
40
60
80
Ni 30Mn 50In 20 100
Ni45Mn35In20
Endothermic
933+771+755
822+660 751+555 911+753
733
731+533
533
711+551
442
400
222
220
331
440
422
Intensity (a.u.)
511+333
(a)
120
2 (deg.)
(b) Ni Mn
Tt
Space group : F d 3 m Prototype : Mn 3Ni 2Si
500
600
B 2/ L 2 1
700
800
Temperature /
In 1000 Fig. 5. (a) XRD pattern taken from the Ni30Mn50In20 alloy, where main peaks are indexed as the Mn3Ni2Si-type structure. (b) A projection of unit cell from [1 0 0] direction for an ideal Mn3Ni2In compound.
(b)
( B 2)
Ni Mn
In
Since the pseudo-boundary surrounding the martensite region ‘m’ appearing in the b (B2) phase region at the annealing temperature, as shown in Fig. 2 corresponds to the region with the martensitic transformation temperature at room temperature, an iso-Ms (= RT) line can be evaluated by EPMA analysis on such a pseudo-boundary between martensite ‘m’ and austenite ‘a’ phase regions, as listed in Table 4 and shown in Fig. 7(a). Due to the thermal history dependence on the microstructure and to the rough boundary structure, no high accuracy of the obtained data can be expected. However, it is seen that the estimated iso-Ms line is almost perfectly consistent with the Ms temperature data determined by DSC for alloy specimens [19] as shown in Fig. 7(b). From Fig. 7(b), the In content of the iso-Ms line can be seen to
|
|
600
| | | | | | | | | | | | | | | | | | ( L 2 1)
|
3.3. Existent regions of martensite and ferromagnetic phases
Te m p e r a t u r e , T
800
by the present authors [18] and the order–disorder transformation 1 temperatures, T B2=L2 , have been determined. In order to obtain t 1 information on T B2=L2 in another section in the b phase, DSC meat surement was performed for some Ni45Mn55 xInx (hereafter 45%Ni) alloys. Fig. 6(a) shows the heating and cooling curves obtained from Ni45Mn35In20 alloy, where endothermic and exothermic reactions corresponding to the B2/L21 order–disorder transformation can be detected at almost the same temperature in the heating and cooling curves, respectively. Here, a temperature with the endothermic peak in the heating curve was defined as the critical 1 temperature T B2=L2 in this second order transformation, as indit cated by the arrow in Fig. 6(a). The obtained data are listed in Table 3 and plotted with open circles in Fig. 6(b), compared with the 1 curve of T B2=L2 estimated from the 50%Ni alloys. It is seen that t 1 the dependence of In content on the T B2=L2 in the 45%Ni section t is basically similar to that in the 50%Ni section, whereas it is lower than that in the 50%Ni section over the whole composition. The 1 maximum temperature (= about 745 °C) of the T B2=L2 is located t at about 23%In as well as that in the 50%Ni section, being about 70 °C lower than that (= 813 °C) in the 50%Ni. From Fig. 6(b), the existent region of the L21 phase can be estimated as shown in the isothermal diagram of 700 °C of Fig. 4(b), where the B2/L21 critical points are indicated by crosses. The existent region of the ordered L21 phase becomes narrow with increasing Mn composition.
400 B 2/ L 2 1
Present work ( T t T 200
10
B 2/ L 2 1 t
with 45at.% Ni)
with 50at.% Ni [18] 20
30
40
In content, x (at.%) Fig. 6. (a) DSC heating and cooling curves of the Ni45Mn35In20 alloy. The critical temperature of order–disorder transition is defined as the temperature with a minimal point in the heating curve as indicated by a bold arrow. (b) Vertical section diagram of 45%Ni in the Ni–Mn–In ternary system, together with that of 50%Ni [18]. The critical compositions of B2/L21 order–disorder transformation at 700 °C are evaluated from this figure as demonstrated with symbol (+).
Table 3 B2/L21 order–disorder transformation temperatures of the 45Ni–Mn–In alloys determined by DSC measurement. Mn
In
Tt
30 25 20 15
15 20 25 30
643 738 734 674
B2=L21
(°C)
decrease with increasing Mn content, and the Ms temperature in the 50%Mn section seems to more drastically decrease with In content than in the 50%Ni section. The ferromagnetic phase region was determined by microstructure observation with magnetic colloid and by EPMA line scanning. Fig. 8(a) shows the OM image in the vicinity of the b phase in the DT specimen annealed at 850 °C that is contrasted by magnetic colloid. It can be seen that a net-like contrast due to magnetic domain walls appears on the b phase layer formed in the diffusion zone. In
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T. Miyamoto et al. / Journal of Alloys and Compounds 549 (2013) 57–63 Table 4 Interface concentration between austenite and martensite phases determined by the diffusion triple method.
(a)
iso-Ms(= RT) (at.%) Mn
In
26.3 31.5 34.8 35.9 46.9 47.2 48.4 49.0 49.2 49.5 50.6 50.7 51.5 52.2 52.0 52.3
21.4 17.0 14.9 14.3 10.6 10.3 10.0 9.6 9.5 9.4 9.3 9.2 9.1 8.9 8.7 8.5
(b)
(a) (c)
(b)
Fig. 8. (a) Microstructure in the vicinity of the b phase in the diffusion triple covered with magnetic colloid and (b) brightness–penetration profile obtained by microstructural analysis along the line indicated in Fig. 8(a). (c) Existing region of ferromagnetic phase at room temperature, where open circles correspond to concentrations determined by chemical analysis for some points on critical boundary between the ferromagnetic and the paramagnetic phases, as demonstrated in Fig. 8(b).
Table 5 Critical concentration of magnetic domain boundary determined by the diffusion triple method. iso-TFM(= RT) (at.%)
Fig. 7. (a) Microstructure in the vicinity of b phase including pseudo-boundary between martensite ‘m’ and austenite ‘a’ phase regions in the diffusion triple, and (b) existing region of martensite phase at room temperature, where open circles correspond to concentrations determined by chemical analysis for some points on boundary between martensite and austenite phases as demonstrated in Fig. 7(a). The data on the Ms temperature [19] obtained from alloy specimens are also shown in Fig. 7(b).
order to determine the ferromagnetic region, image analyses were carried out along some lines across to the ferromagnetic layer, as shown in Fig. 8(b). Since the baseline of contrast, except for the ferromagnetic region with a dark contrast, gradually decreases to the right, the critical boundary between ferromagnetic and paramagnetic regions, iso-TFM (= RT), can be determined by chemical
Mn
Ni
49.6 52.1 48.6 46.1 46.2 25.8 26.7 42.0 40.1 39.3 21.4 35.2 40.4
13.3 9.5 10.6 10.9 10.7 29.4 28.5 12.0 13.4 13.3 30.3 14.5 13.3
analysis on the boundary where the magnetic domain contrast disappears. The concentrations at the critical points for some different scanning lines in the b phase are listed in Table 5 and plotted with
T. Miyamoto et al. / Journal of Alloys and Compounds 549 (2013) 57–63
open circles in the partial phase diagram of 800 °C in Fig. 8(c). It is seen that a couple of iso-TFM lines are located in the vicinities of about 10–20 and 25–30 at.% In. Thus, the DT technique is very effective for rough estimation of the existent region of the ferromagnetic phase as well as the martensite phase. It is interesting to note that the iso-Ms line drawn in Fig. 7(b) perfectly coincides with the iso-TFM line appearing in the lower In region. As mentioned in Section 1, the martensitic transformation in the Ni-rich NiMnIn alloy is from the ferromagnetic austenite to the paramagnetic martensite phase [3,20]. From the present results, it is expected that this unique transformation occurs in the wide composition range from the Ni-rich region to the Mn-rich region, which is coincident with a previous report by Llamazares et al. [21]. On the other hand, the iso-TFM line at about 28–29%In corresponds to the critical compositions of the normal ferromagnetic transition in the austenite phase, which is in good agreement to the Curie temperature of the slightly In-rich Ni50Mn24In26 alloy [22]. 4. Conclusions The phase equilibria of 700 and 850 °C in the Ni–Mn–In system were determined using diffusion triples and multi-phase alloys. It was confirmed that a single-phase region of the b phase exists in a wide composition range parallel to the Mn–In section and that the martensite region observed at room temperature in the b phase region is also located over a wide range with 0–20 at.% In. Moreover, the ferromagnetic region at room temperature in the b phase was estimated by image analysis of the microstructure contrasted by magnetic colloid for the DT specimen, the result being in good agreement with the previous data by the conventional alloy method. Acknowledgments The authors would like to thank Drs. R.Y. Umetsu and W. Ito for their help in the experimental work. This work was supported by a
63
Grant-in-Aid for Scientific Research (S) from the Japan Society for the Promotion of Science (JSPS). References [1] Y. Sutou, Y. Imano, N. Koeda, T. Omori, R. Kainuma, K. Ishida, K. Oikawa, Appl. Phys. Lett. 85 (2004) 4358. [2] T. Krenke, E. Duman, M. Acet, E.F. Wassermann, X. Moya, L. Manosa, A. Planes, Nat. Mater. 4 (2005) 450. [3] K. Oikawa, W. Ito, Y. Imano, Y. Sutou, R. Kainuma, K. Ishida, Appl. Phys. Lett. 88 (2006) 122507. [4] R. Kainuma, Y. Imano, W. Ito, Y. Sutou, H. Morito, S. Okamoto, O. Kitakami, K. Oikawa, A. Fujita, T. Kanomata, K. Ishida, Nature 439 (2006) 957. [5] R. Kainuma, Y. Imano, W. Ito, H. Morito, Y. Sutou, K. Oikawa, A. Fujita, K. Ishida, Appl. Phys. Lett. 88 (2006) 192513. [6] Z.D. Han, D.H. Wang, C.L. Zhang, S.L. Tang, B.X. Gu, Y.W. Du, Appl. Phys. Lett. 89 (2006) 182507. [7] K. Koyama, H. Okada, K. Watanabe, T. Kanomata, R. Kainuma, W. Ito, K. Oikawa, K. Ishida, Appl. Phys. Lett. 89 (2006) 182510. [8] S.Y. Yu, Z.H. Liu, G.D. Liu, J.L. Chen, Z.X. Cao, G.H. Wu, B. Zhang, X.X. Zhang, Appl. Phys. Lett. 89 (2006) 162503. [9] B. Zhang, X.X. Zhang, S.Y. Yu, J.L. Chen, Z.X. Cao, G.H. Wu, Appl. Phys. Lett. 91 (2007) 012510. [10] A. Planes, L. Manosa, M. Acet, J. Phys. 21 (2009) 233201. [11] T.B. Massalski, in: T.B. Massalski (Ed.), Binary Alloys Phase Diagrams, Second ed., ASM International, 1990, p. 2580. [12] C. Guo, X. Du, Intermetallics 13 (2005) 525. [13] T.B. Massalski, in: T.B. Massalski (Ed.), Binary Alloys Phase Diagrams, Second ed., ASM International, 1990, p. 2259. [14] T.B. Massalski, in: T.B. Massalski (Ed.), Binary Alloys Phase Diagrams, Second ed., ASM International, 1990, p. 2267. [15] Z. Jin, Scand. J. Metall. 10 (1981) 279. [16] J.C. Zhao, in: J.C. Zhao (Ed.), Methods for Phase Diagram Determination, Elsevier, Oxford, 2007, p. 246. [17] R. Ducher, R. Kainuma, I. Ohnuma, K. Ishida, J. Alloys Comp. 437 (2007) 93– 101. [18] T. Miyamoto, W. Ito, R.Y. Umetsu, R. Kainuma, T. Kanomata, K. Ishida, Scripta Mater. 62 (2010) 151. [19] W. Ito, R. Kainuma, unpublished work. [20] V.V. Khovaylo, T. Kanomata, T. Tanaka, M. Nakashima, Y. Amako, R. Kainuma, R.Y. Umetsu, H. Morito, H. Miki, Phys. Rev. 80 (2009) 144409. [21] J.L. Llamazares, T. Sanchez, J.D. Santos, M.J. Perez, M.L. Sanchez, B. Hernando, L. Escoda, J.J. Sunol, R. Varga, Appl. Phys. Lett. 92 (2008) 012513. [22] T. Miyamoto, W. Ito, R.Y. Umetsu, T. Kanomata, K. Ishida, R. Kainuma, Mater. Trans. 52 (2011) 1836–1839.
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Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Phase equilibria in the Ni–Mn–In alloy system T. Miyamoto, M. Nagasako ⇑, R. Kainuma Department of Materials Science, Graduate School of Engineering, Tohoku University, 6-6-02 Aoba, Sendai 980-8579, Japan
a r t i c l e
i n f o
Article history: Received 12 July 2012 Received in revised form 29 August 2012 Accepted 30 August 2012 Available online 6 September 2012 Keywords: Phase equilibria Diffusion triple Ni–Mn–In Martensitic transformation Order–disorder transition Magnetic transition
a b s t r a c t Phase equilibria at 700 and 850 °C, critical temperatures of B2/L21 order–disorder transformation and martensitic and ferromagnetic phase regions at room temperature in the Ni–Mn–In system were determined mainly by diffusion triple method using a two-stage diffusion couple technique. It was confirmed that a single phase region of the b phase at both 700 and 850 °C exists in a wide composition range along the NiMn–NiIn section and that the L21 ordered phase region appears in the vicinity of Ni2MnIn in the temperature region below about 800 °C. The composition lines, iso-Ms and iso-TFM, possessing the Ms and TC temperatures at room temperature, respectively, were successfully estimated and the coincidence between the iso-Ms and iso-TFM was confirmed in the composition region from 10 to 20 at.% In in the bphase region. Ó 2012 Elsevier B.V. All rights reserved.
1. Introduction Since unusual martensitic transformation from the ferromagnetic austenite phase to the paramagnetic martensite phase has been reported in the off-stoichiometric Ni2MnIn and Ni2MnSn alloys with Heusler L21-type ordered bcc structure [1–4], Ni–Mn–In based alloys have received much attention as high-performance multifunctional materials. The martensitic transformation temperatures of these alloys are drastically decreased by application of a magnetic field and magnetic-field-induced reverse transformation, namely, metamagnetic phase transition, occurs near the martensitic reverse transformation start temperature, Ms [2,3]. Moreover, a magnetic-field-induced shape memory (SM) effect, i.e., metamagnetic SM effect, was confirmed in Ni45Co5Mn36.7In13.3 [4] and Ni43Co7Mn39Sn11 [5] alloys at room temperature, and many interesting phenomena accompanying this unique transformation, such as the inverse magnetocaloric effect [2,6], the giant magnetoresistance effect [7,8] and giant magnetothermal conductivity [9] were reported. While the basic physical properties of the Ni–Mn–In Heusler alloys have been reviewed by Planes et al. [10], no phase diagram covering the whole range of the Ni–Mn–In ternary system has been established. Fig. 1 shows the most reliable Ni–Mn [11,12], Mn–In [13] and Ni–In [14] binary phase diagrams to date. It is seen that the single phase region of the a (A1: disordered fcc Ni or Mn) exists in a wide composition range in the Ni–Mn binary system and that the b (B2: CsCl-type ordered bcc NiMn or NiIn) phase in relation to the
⇑ Corresponding author. Tel./fax: +81 22 795 5256. E-mail address: [email protected] (M. Nagasako). 0925-8388/$ - see front matter Ó 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2012.08.128
Ni2MnIn Heusler alloy appears in the temperature range from 700 to 1000 °C in the vicinity of the stoichiometric composition in the Ni–Mn and Ni–In systems. From this information, the b-NiMn phase region is expected to be connected to the b-NiIn phase region at high temperatures. In spite of the importance of further alloy development, only limited information on the phase equilibria at elevated temperatures in this ternary system has been reported. In the present study, the phase equilibria at 700 and 850 °C in whole region of the Ni–Mn–In ternary system are determined and the composition region where martensite and ferromagnetic phases are observed at room temperature is evaluated by using the diffusion triple (DT) method. Here, the DT method is a kind of combinatorial methods, first reported by Jin [15] and more recently employed by Zhao et al. [16]. 2. Experimental procedures 2.1. Diffusion triples Details of the DT method that we used in the present study were reported in our previous paper on the Co–Fe–Ga phase diagram [17]. The DT specimens for the Ni– Mn–In system were prepared by the following three steps: (1) Pure Ni (99.9%) and Mn (99.9%) blocks were joined with induction melting in an argon atmosphere, and only the Mn block laid on a Ni block was melted in the induction furnace. Ni/Mn diffusion couple (DC) samples were diffusion-annealed at 1000 °C for 168 h in evacuated quartz tubes to obtain a continuous diffusion zone between the Ni and Mn blocks. (2) A cylindrical hole with a diameter of 3 mm was formed by electro-spark machining near the diffusion zone of the Ni/Mn DC, and some In (99.99%) chips were inserted into the hole after removal of the surface layer by polishing. Diffusion triples were finally obtained by diffusion-annealing at 700 °C for 24 h, 96 or 168 h and at 850 °C for 3, 12, or 24 h in evacuated quartz tubes.
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Fig. 1. Binary phase diagrams constituting the Ni–Mn–In ternary system [11–14].
Table 1 Equilibrium concentration in the Ni–Mn–In system determined by the diffusion triple method. Temp. (°C)
Time (hours)
Phase U
Phase C
Equilibrium composition (at.%) Phase U
850
3
f
c b 12
c c c c b
c 24
b-Mn
c b-Mn 700
24
c Ni2In3
c Mn3Ni2In b
c
96 168
b Mn3Ni2In b b b
c c f
c
Ni3In b L b f f b f b-Mn L b-Mn L b-Mn L b b-Mn Mn3Ni2In b-Mn Mn3Ni2In L Mn3Ni2In Ni2In3 Ni2In3 b f Ni3In Ni3In
Phase C
Mn
In
Mn
In
0.4 22.1 2.9 41.5 15.0 10.9 28.0 12.5 4.0 68.8 67.2 64.6 72.1 1.6 67.3 56.7 32.0 69.9 42.2 48.0 51.4 9.3 7.7 20.8 13.2 1.4 9.4
30.9 7.6 57.2 0.3 6.2 6.6 2.8 32.2 34.2 18.0 2.4 11.0 0.8 63.3 1.2 13.3 28.8 1.3 15.6 21.3 7.8 46.3 48.5 3.3 4.2 30.4 2.8
0.3 20.5 6.6 45.6 5.7 3.5 25.9 4.8 10.8 56.3 65.5 54.1 71.9 0.8 55.8 69.1 48.3 69.2 49.3 36.6 53.4 3.0 3.6 19.2 3.2 1.4 1.8
24.6 25.4 62.7 4.1 30.6 30.3 21.8 32.6 33.8 27.9 6.0 20.4 4.4 85.5 6.1 5.1 21.1 5.0 19.7 40.4 15.1 55.3 55.0 24.9 30.5 24.6 24.2
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T. Miyamoto et al. / Journal of Alloys and Compounds 549 (2013) 57–63 Table 2 Equilibrium concentration in the Ni–Mn–In system determined by the conventional alloy method. Composition (at.%)
Temp. (°C)
Time (hours)
Phase U
Phase C
Phase H
Equilibrium composition (at.%) Phase U
Mn
In
78 59 55
17 6 40 30 5 10 32 16 28 38
850
17 6 40 30 20 25 18 5 32 16 15 50 51 32 28 45 38
700
41 36 30 8 4 78 59 55
50 42 41 30 25 15 12 8 5 4
24
b-Mn
c Mn3In b-Mn
c c b
c c b 72
b-Mn
c Mn3In Mn3In Mn3Ni2In Mn3Ni2In b
c b
c c b b b
c b b
Phase C
Phase H Mn
In
Mn3In b L L b b L b b f
– L – – – – – – f –
79.8 65.7 70.2 69.8 38.6 34.5 29.0 31.1 17 9.3
15.3 2.4 26.0 16.7 0.9 1.7 26.4 2.5 6.3 37.9
70.6 55.9 47.6 54.2 41.1 35.4 26.8 29.9 6.1 3.1
25.1 8.0 45.3 29.7 6.6 11.6 43.4 17.4 30.3 35.9
– 56.3 – – – – – – 13.6 –
– 17.7 – – – – – – 28.1 –
Mn3In b L L b-Mn Mn3In Mn3Ni2In b Mn3Ni2In b b L L f b NiIn f
– Mn3Ni2In – – – L – b’ L – – – – – f – Ni13In9
79.1 67.9 65.6 58.2 52.0 51.0 40.7 42.3 29.6 31.8 28.5 16.8 13.4 15.5 17.6 10.0 11.7
15.8 1.4 28.2 26.5 20.4 22.4 16.9 1.1 30.9 1.1 1.6 41.8 45.0 29.1 4.3 40.0 35.9
73.6 53.0 30.2 36.2 64.9 57.9 49.2 42.6 45.0 29.2 25.8 27.7 6.9 5.4 17.3 0.3 2.7
21.7 7.7 62.8 42.6 16.4 25.7 18.5 6.6 23.5 19.5 22.0 50.3 73.8 31.5 25.8 48.9 35.5
– 54.3 – – – 35.9 – 39.4 32.2 – – – – – 5.0 – 1.7
– 14.9 – – – 42.0 – 0.6 42.8 – – – – – 30.3 – 39.3
(3) The obtained DTs were cut from a section vertical to the In cylinder and subsequent to polishing were chemically analyzed using an electron microprobe analyzer (EPMA) along some lines parallel to the Ni/Mn interface. Consequently, information on phase equilibria over a wide composition range was extracted from the line analyses along the diffusion paths, as listed in Table 1.
Mn
In
Mn
In
(a)
In the present study, the obtained DTs covered most of the composition range in the ternary system except for the In-rich corner in which a liquid phase appeared. 2.2. Alloy specimens The alloy samples shown in Table 2 were additionally prepared and the microstructures were analyzed by EPMA in order to confirm reliability of the phase equilibrium determined by the DTs. Ni30Mn50In20 alloy was also prepared to determine the crystal structure of a new phase detected in the DT examination. These alloys were induction-melted from pure Ni, Mn, and In, and annealed for equilibration at 700 °C for 72 h and at 850 °C for 24 h in quartz tubes under an argon atmosphere. Microstructure observation was carried out by optical microscopy (OM). The crystal structure was determined by transmission electron microscopic (TEM) examination for thin foils and by X-ray powder diffraction (XRD), where the TEM sample was prepared by jet-electropolishing with a solution of 8 vol.% perchloric acid, 72 vol.% acetic acid, 8 vol.% ethylene glycol and 12 vol.% ethanol at room temperature. The order–disorder transformation temperature was determined by differential scanning calorimetry (DSC) at a heating rate of 10 °C/min.
(b)
2.3. Determination of iso-Ms and -TFM lines Regions of alloy composition possessing martensite and ferromagnetic phases at room temperature were determined from the microstructure in the DTs. In the case of the martensite phase, since the chemical compositions of the alloys possessing the martensitic transformation starting temperature T Ms at room temperature correspond to those at the edge of the martensite region in the DTs, the iso-Ms (= RT) composition line can be estimated by chemical analysis at many positions on the edge of the martensite region, as reported in our previous papers [17]. It is impossible to distinguish between ferromagnetic and non-magnetic regions by conventional OM observation. In the present study, the ferromagnetic region was determined by contrast of magnetic walls visualized by magnetic colloids.
Fig. 2. (a) Microstructure and (b) concentration–penetration profiles for Ni and Mn elements in the vicinity of the diffusion zone in the diffusion triple annealed at 850 °C. The martensite region partially formed in the b (B2) phase region during cooling from annealing temperature is labeled with ‘m’. Equilibrium concentrations were evaluated by extrapolation of the profiles to the phase interface.
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The chemical compositions of the alloys possessing the ferromagnetic to nonmagnetic transition temperature at room temperature, i.e., the iso-TFM (= RT) compositions, were evaluated by chemical analysis at many positions on the edge of the ferromagnetic phase region as well as the iso-Ms (= RT) compositions. Here, the edge of the ferromagnetic phase region was determined by using image analysis.
(a)
3. Results and discussion 3.1. Isothermal section diagrams of 700 and 850 °C Fig. 2(a) shows the OM image in the vicinity of diffusion zone in the DT specimen annealed at 850 °C, where three distinct regions, i.e., Ni-rich c, Mn-rich c and In-rich b phases, are observed on the left, lower right and center sides. Besides these phases, a martensite phase region with a lamella-like contrast and a liquid phase region with a solidification structure are observed as labeled with ‘m’ and ‘Liquid’ in Fig. 2(a), respectively, where the martensite phase region is formed in the b phase region and there is a phase boundary with the austenite phase region with no lamella-like contrast. Since the martensitic transformation occurs during cooling, the phase boundary dividing the austenite from the martensite region corresponds to a habit plane and has no relation to the phase diagram at the annealing temperature. Equilibrium compositions among those observed phases were determined using concentration–penetration profiles obtained by EPMA for the DTs annealed at 700 and 850 °C. Fig. 2(b) shows typical concentration–penetration profiles for Mn and Ni elements in the vicinity of the diffusion zone in the DT specimen annealed at 850 °C. Phase equilibrium concentrations were determined by extrapolation of the concentration profiles to the phase interface as demonstrated in Fig. 2(b). All the results obtained from the DT samples annealed at 700 and 850 °C are listed in Table 1. Because it is difficult to determine three-phase and solid + liquid two-phase equilibria using the DT method, some alloy samples listed in Table 2 were also prepared. A typical microstructure obtained from the Ni35Mn59In6 alloy is shown in Fig. 3. In this specimen annealed at 850 °C followed by quenching, three phases, the c, the b and the Liquid (L) phases, are observed, whereas the b phase has transformed to the martensite phase with the lamellalike contrast during cooling from the annealing temperature. Average phase equilibrium concentrations were determined from EPMA data obtained from at least seven different points for each phase. All the results obtained from the DT samples and alloy specimens annealed at 850 and 700 °C are plotted in Fig. 4(a) and (b), respectively. It is clear that phase equilibria determined by the DT method are coincident with those by the alloy method. In the isothermal diagrams of 700 and 850 °C, the topological constitutions on the phase equilibria are not complicated and are similar to each other, except for the presence of Mn3Ni2In compound at 700 °C, where the Mn3Ni2In phase is a ternary intermetallic com-
(b)
Fig. 4. Phase equilibria at (a) 850 °C and (b) 700 °C of Ni–Mn–In ternary system determined with diffusion triples and alloy specimens. Symbol (+) plotted in b phase of 700 °C diagram indicates points on critical boundary for the chemical 1 ordering estimated from the T B2=L2 curves of Fig. 6. t
pound that is firstly found in the present study. The liquid phase widely extends to the middle region of the ternary diagrams and the NiMn phase with B2 structure exists along the NiMn–NiIn section including the Heusler composition Ni2MnIn. According to the Ni–In binary phase diagram shown Fig. 1 [14], three binary compounds Ni13In9 (B81: NiAs-type), NiIn (B35: CoSn-type), and Ni2In3 (D513: Al3Ni2-type) exist up to 876, 860 and 870 °C, respectively. All these phases, however, do not appear in the 850 °C diagram. This means that the phase equilibria in the Ni–In binary system should be more accurately reexamined. On the other hand, all these binary phases appear in the 700 °C diagram, while their existing regions are limited on the binary side. The basic crystal structure of the Mn3Ni2In compound was determined by TEM observation. From analysis of obtained electron diffraction patterns, the crystal structure of this phase was determined to belong to the cubic system with space group Fig. 5(a) shows the XRD pattern taken from the Ni30Mn50Fd 3m. In20 single-phase alloy. Based on the space group and the experimental XRD pattern, the prototype of this phase is conformed as being the Mn3Ni2Si structure in which Mn, Ni, and Si (In) atoms occupy 48f, 32e, 16d sites, respectively, as shown in Fig. 5(b). The lattice constant of this phase in the Ni30Mn50In20 alloy is estimated as a = 1.1307 nm. 3.2. Order–disorder transformation in the Ni(Mn, In) phase
Fig. 3. Three-phase microstructure taken from Ni35Mn59In6 alloy annealed at 850 °C for 24 h, where the b (B2) phase has transformed to the martensite phase during quenching.
Order–disorder transformation from the B2 to L21 phase in Ni50Mn50 xInx (hereafter 50%Ni) alloys has recently been investigated
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(a)
20
40
60
80
Ni 30Mn 50In 20 100
Ni45Mn35In20
Endothermic
933+771+755
822+660 751+555 911+753
733
731+533
533
711+551
442
400
222
220
331
440
422
Intensity (a.u.)
511+333
(a)
120
2 (deg.)
(b) Ni Mn
Tt
Space group : F d 3 m Prototype : Mn 3Ni 2Si
500
600
B 2/ L 2 1
700
800
Temperature /
In 1000 Fig. 5. (a) XRD pattern taken from the Ni30Mn50In20 alloy, where main peaks are indexed as the Mn3Ni2Si-type structure. (b) A projection of unit cell from [1 0 0] direction for an ideal Mn3Ni2In compound.
(b)
( B 2)
Ni Mn
In
Since the pseudo-boundary surrounding the martensite region ‘m’ appearing in the b (B2) phase region at the annealing temperature, as shown in Fig. 2 corresponds to the region with the martensitic transformation temperature at room temperature, an iso-Ms (= RT) line can be evaluated by EPMA analysis on such a pseudo-boundary between martensite ‘m’ and austenite ‘a’ phase regions, as listed in Table 4 and shown in Fig. 7(a). Due to the thermal history dependence on the microstructure and to the rough boundary structure, no high accuracy of the obtained data can be expected. However, it is seen that the estimated iso-Ms line is almost perfectly consistent with the Ms temperature data determined by DSC for alloy specimens [19] as shown in Fig. 7(b). From Fig. 7(b), the In content of the iso-Ms line can be seen to
|
|
600
| | | | | | | | | | | | | | | | | | ( L 2 1)
|
3.3. Existent regions of martensite and ferromagnetic phases
Te m p e r a t u r e , T
800
by the present authors [18] and the order–disorder transformation 1 temperatures, T B2=L2 , have been determined. In order to obtain t 1 information on T B2=L2 in another section in the b phase, DSC meat surement was performed for some Ni45Mn55 xInx (hereafter 45%Ni) alloys. Fig. 6(a) shows the heating and cooling curves obtained from Ni45Mn35In20 alloy, where endothermic and exothermic reactions corresponding to the B2/L21 order–disorder transformation can be detected at almost the same temperature in the heating and cooling curves, respectively. Here, a temperature with the endothermic peak in the heating curve was defined as the critical 1 temperature T B2=L2 in this second order transformation, as indit cated by the arrow in Fig. 6(a). The obtained data are listed in Table 3 and plotted with open circles in Fig. 6(b), compared with the 1 curve of T B2=L2 estimated from the 50%Ni alloys. It is seen that t 1 the dependence of In content on the T B2=L2 in the 45%Ni section t is basically similar to that in the 50%Ni section, whereas it is lower than that in the 50%Ni section over the whole composition. The 1 maximum temperature (= about 745 °C) of the T B2=L2 is located t at about 23%In as well as that in the 50%Ni section, being about 70 °C lower than that (= 813 °C) in the 50%Ni. From Fig. 6(b), the existent region of the L21 phase can be estimated as shown in the isothermal diagram of 700 °C of Fig. 4(b), where the B2/L21 critical points are indicated by crosses. The existent region of the ordered L21 phase becomes narrow with increasing Mn composition.
400 B 2/ L 2 1
Present work ( T t T 200
10
B 2/ L 2 1 t
with 45at.% Ni)
with 50at.% Ni [18] 20
30
40
In content, x (at.%) Fig. 6. (a) DSC heating and cooling curves of the Ni45Mn35In20 alloy. The critical temperature of order–disorder transition is defined as the temperature with a minimal point in the heating curve as indicated by a bold arrow. (b) Vertical section diagram of 45%Ni in the Ni–Mn–In ternary system, together with that of 50%Ni [18]. The critical compositions of B2/L21 order–disorder transformation at 700 °C are evaluated from this figure as demonstrated with symbol (+).
Table 3 B2/L21 order–disorder transformation temperatures of the 45Ni–Mn–In alloys determined by DSC measurement. Mn
In
Tt
30 25 20 15
15 20 25 30
643 738 734 674
B2=L21
(°C)
decrease with increasing Mn content, and the Ms temperature in the 50%Mn section seems to more drastically decrease with In content than in the 50%Ni section. The ferromagnetic phase region was determined by microstructure observation with magnetic colloid and by EPMA line scanning. Fig. 8(a) shows the OM image in the vicinity of the b phase in the DT specimen annealed at 850 °C that is contrasted by magnetic colloid. It can be seen that a net-like contrast due to magnetic domain walls appears on the b phase layer formed in the diffusion zone. In
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T. Miyamoto et al. / Journal of Alloys and Compounds 549 (2013) 57–63 Table 4 Interface concentration between austenite and martensite phases determined by the diffusion triple method.
(a)
iso-Ms(= RT) (at.%) Mn
In
26.3 31.5 34.8 35.9 46.9 47.2 48.4 49.0 49.2 49.5 50.6 50.7 51.5 52.2 52.0 52.3
21.4 17.0 14.9 14.3 10.6 10.3 10.0 9.6 9.5 9.4 9.3 9.2 9.1 8.9 8.7 8.5
(b)
(a) (c)
(b)
Fig. 8. (a) Microstructure in the vicinity of the b phase in the diffusion triple covered with magnetic colloid and (b) brightness–penetration profile obtained by microstructural analysis along the line indicated in Fig. 8(a). (c) Existing region of ferromagnetic phase at room temperature, where open circles correspond to concentrations determined by chemical analysis for some points on critical boundary between the ferromagnetic and the paramagnetic phases, as demonstrated in Fig. 8(b).
Table 5 Critical concentration of magnetic domain boundary determined by the diffusion triple method. iso-TFM(= RT) (at.%)
Fig. 7. (a) Microstructure in the vicinity of b phase including pseudo-boundary between martensite ‘m’ and austenite ‘a’ phase regions in the diffusion triple, and (b) existing region of martensite phase at room temperature, where open circles correspond to concentrations determined by chemical analysis for some points on boundary between martensite and austenite phases as demonstrated in Fig. 7(a). The data on the Ms temperature [19] obtained from alloy specimens are also shown in Fig. 7(b).
order to determine the ferromagnetic region, image analyses were carried out along some lines across to the ferromagnetic layer, as shown in Fig. 8(b). Since the baseline of contrast, except for the ferromagnetic region with a dark contrast, gradually decreases to the right, the critical boundary between ferromagnetic and paramagnetic regions, iso-TFM (= RT), can be determined by chemical
Mn
Ni
49.6 52.1 48.6 46.1 46.2 25.8 26.7 42.0 40.1 39.3 21.4 35.2 40.4
13.3 9.5 10.6 10.9 10.7 29.4 28.5 12.0 13.4 13.3 30.3 14.5 13.3
analysis on the boundary where the magnetic domain contrast disappears. The concentrations at the critical points for some different scanning lines in the b phase are listed in Table 5 and plotted with
T. Miyamoto et al. / Journal of Alloys and Compounds 549 (2013) 57–63
open circles in the partial phase diagram of 800 °C in Fig. 8(c). It is seen that a couple of iso-TFM lines are located in the vicinities of about 10–20 and 25–30 at.% In. Thus, the DT technique is very effective for rough estimation of the existent region of the ferromagnetic phase as well as the martensite phase. It is interesting to note that the iso-Ms line drawn in Fig. 7(b) perfectly coincides with the iso-TFM line appearing in the lower In region. As mentioned in Section 1, the martensitic transformation in the Ni-rich NiMnIn alloy is from the ferromagnetic austenite to the paramagnetic martensite phase [3,20]. From the present results, it is expected that this unique transformation occurs in the wide composition range from the Ni-rich region to the Mn-rich region, which is coincident with a previous report by Llamazares et al. [21]. On the other hand, the iso-TFM line at about 28–29%In corresponds to the critical compositions of the normal ferromagnetic transition in the austenite phase, which is in good agreement to the Curie temperature of the slightly In-rich Ni50Mn24In26 alloy [22]. 4. Conclusions The phase equilibria of 700 and 850 °C in the Ni–Mn–In system were determined using diffusion triples and multi-phase alloys. It was confirmed that a single-phase region of the b phase exists in a wide composition range parallel to the Mn–In section and that the martensite region observed at room temperature in the b phase region is also located over a wide range with 0–20 at.% In. Moreover, the ferromagnetic region at room temperature in the b phase was estimated by image analysis of the microstructure contrasted by magnetic colloid for the DT specimen, the result being in good agreement with the previous data by the conventional alloy method. Acknowledgments The authors would like to thank Drs. R.Y. Umetsu and W. Ito for their help in the experimental work. This work was supported by a
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Grant-in-Aid for Scientific Research (S) from the Japan Society for the Promotion of Science (JSPS). References [1] Y. Sutou, Y. Imano, N. Koeda, T. Omori, R. Kainuma, K. Ishida, K. Oikawa, Appl. Phys. Lett. 85 (2004) 4358. [2] T. Krenke, E. Duman, M. Acet, E.F. Wassermann, X. Moya, L. Manosa, A. Planes, Nat. Mater. 4 (2005) 450. [3] K. Oikawa, W. Ito, Y. Imano, Y. Sutou, R. Kainuma, K. Ishida, Appl. Phys. Lett. 88 (2006) 122507. [4] R. Kainuma, Y. Imano, W. Ito, Y. Sutou, H. Morito, S. Okamoto, O. Kitakami, K. Oikawa, A. Fujita, T. Kanomata, K. Ishida, Nature 439 (2006) 957. [5] R. Kainuma, Y. Imano, W. Ito, H. Morito, Y. Sutou, K. Oikawa, A. Fujita, K. Ishida, Appl. Phys. Lett. 88 (2006) 192513. [6] Z.D. Han, D.H. Wang, C.L. Zhang, S.L. Tang, B.X. Gu, Y.W. Du, Appl. Phys. Lett. 89 (2006) 182507. [7] K. Koyama, H. Okada, K. Watanabe, T. Kanomata, R. Kainuma, W. Ito, K. Oikawa, K. Ishida, Appl. Phys. Lett. 89 (2006) 182510. [8] S.Y. Yu, Z.H. Liu, G.D. Liu, J.L. Chen, Z.X. Cao, G.H. Wu, B. Zhang, X.X. Zhang, Appl. Phys. Lett. 89 (2006) 162503. [9] B. Zhang, X.X. Zhang, S.Y. Yu, J.L. Chen, Z.X. Cao, G.H. Wu, Appl. Phys. Lett. 91 (2007) 012510. [10] A. Planes, L. Manosa, M. Acet, J. Phys. 21 (2009) 233201. [11] T.B. Massalski, in: T.B. Massalski (Ed.), Binary Alloys Phase Diagrams, Second ed., ASM International, 1990, p. 2580. [12] C. Guo, X. Du, Intermetallics 13 (2005) 525. [13] T.B. Massalski, in: T.B. Massalski (Ed.), Binary Alloys Phase Diagrams, Second ed., ASM International, 1990, p. 2259. [14] T.B. Massalski, in: T.B. Massalski (Ed.), Binary Alloys Phase Diagrams, Second ed., ASM International, 1990, p. 2267. [15] Z. Jin, Scand. J. Metall. 10 (1981) 279. [16] J.C. Zhao, in: J.C. Zhao (Ed.), Methods for Phase Diagram Determination, Elsevier, Oxford, 2007, p. 246. [17] R. Ducher, R. Kainuma, I. Ohnuma, K. Ishida, J. Alloys Comp. 437 (2007) 93– 101. [18] T. Miyamoto, W. Ito, R.Y. Umetsu, R. Kainuma, T. Kanomata, K. Ishida, Scripta Mater. 62 (2010) 151. [19] W. Ito, R. Kainuma, unpublished work. [20] V.V. Khovaylo, T. Kanomata, T. Tanaka, M. Nakashima, Y. Amako, R. Kainuma, R.Y. Umetsu, H. Morito, H. Miki, Phys. Rev. 80 (2009) 144409. [21] J.L. Llamazares, T. Sanchez, J.D. Santos, M.J. Perez, M.L. Sanchez, B. Hernando, L. Escoda, J.J. Sunol, R. Varga, Appl. Phys. Lett. 92 (2008) 012513. [22] T. Miyamoto, W. Ito, R.Y. Umetsu, T. Kanomata, K. Ishida, R. Kainuma, Mater. Trans. 52 (2011) 1836–1839.