Metallurgical and Materials Transactions A

, 42:3062

Structural Transitions and Magnetic Properties of Ni50Mn36.7In13.3 Particles with Amorphous-Like Phase

Authors

  • D. M. Liu
    • Key Laboratory for Anisotropy and Texture of Materials (Ministry of Education)Northeastern University
  • Z. H. Nie
    • School of Materials Science and Engineering, Beijing Institute of Technology
  • Y. Ren
    • X-ray Science DivisionArgonne National Laboratory
    • School of Materials Science and Engineering, Beijing Institute of Technology
  • J. Pearson
    • Materials Science DivisionArgonne National Laboratory
  • P. K. Liaw
    • Department of Materials Science and EngineeringThe University of Tennessee
  • D. E. Brown
    • Department of PhysicsNorthern Illinois University
Article

DOI: 10.1007/s11661-011-0724-x

Cite this article as:
Liu, D.M., Nie, Z.H., Ren, Y. et al. Metall and Mat Trans A (2011) 42: 3062. doi:10.1007/s11661-011-0724-x
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Abstract

The high-energy ball-milling method was used for fabricating Ni50Mn36.7In13.3 fine-sized particles. The as-melt polycrystalline Ni50Mn36.7In13.3 alloy exhibits a 14 M modulated martensite structure at room temperature (RT). The atomic pair distribution function analysis together with the differential scanning calorimetry technique proved that the 14 M modulated martensite transformed to a metastable amorphous-like structure after ball milling for 8 hours. Annealing of the ball-milled particles with the amorphous-like phase first led to the crystallization to form a B2 structure at 523 K (250 °C), and then an ordered Heusler L21 structure (with a small tetragonal distortion) at 684 K (411 °C). The annealed particles undergo different structural transitions during cooling, tailored by the atomic arrangements of the high-temperature phase. Low-field thermomagnetization measurements show that the ball-milled particles with the amorphous-like structure or the atomically disordered crystalline structure exhibit a magnetic transition from the paramagnetic-like to the spin-glass state with decreasing temperature, whereas the crystalline particles with the ordered Heusler L21 structure present a ferromagnetic behavior with the Curie temperature Tc ≈ 310 K (37 °C).

1 Introduction

The Ni-Mn-Ga/Ni-Mn-In ferromagnetic shape memory alloys (FSMAs) have attracted much interest due to the colossal magnetic-field-induced strain (MFIS). The MFIS can be achieved through the reorientation of martensitic variants (e.g., Ni-Mn-Ga[13]) or reverse martensitic phase transformation (e.g., Ni-Mn-In[48]) while a magnetic field is applied. Besides the colossal MFIS, the Ni-Mn-In FSMAs also exhibit the giant magnetoresistance and large inverse magnetocaloric effect.[711] The phase transformation behaviors and structural characteristics of the Ni-Mn-Ga/Ni-Mn-In FSMAs have been well documented.[111] The previous studies focused mainly on bulk polycrystalline materials and single crystals. However, the intrinsic brittleness of bulk polycrystalline alloys and high cost of single crystals have significantly hindered their practical applications.

To overcome this problem, and driven by the motivation to design new types of functional materials used as nanoscale actuators, Wang and co-workers[12,13] have prepared Ni-Mn-Ga nanoparticles through the high-energy ball-milling method. It is found that the Ni-Mn-Ga nanoparticles prepared by ball milling showed a disordered face-centered-cubic structure, which transformed back to the ordered Heusler structure upon annealing. The annealed Ni-Mn-Ga nanoparticles undergo new sequences of martensitic structural transitions, different from their coarse-grained counterparts.[12,13] Actually, high-energy ball milling provides an effective routine in synthesizing a variety of alloy phases with the equilibrium or nonequilibrium state. For different intermetallic compounds with a variety of ordered atomic arrangements, the ball-milling process can induce amorphous particles, nanostructured particles, metastable crystalline, quasi-crystalline particles, etc.[1222] Zhou and Bakker[18,19] reported that both ordered Co2Ge and GdAl2 are ferromagnets at lower temperatures. After ball milling, Co2Ge eventually transforms to the amorphous state,[18] while GdAl2 disorders atomically but remains crystalline.[19] The milling intensities can also affect the nature of the particles.[17,20] It is reported that both atomically disordered crystalline CoZr and amorphous CoZr particles can be obtained by ball milling, depending on the milling time.[20] Moreover, ball milling may induce particles with unique magnetic properties due to the deformation-induced atomic rearrangements. It is found that both the amorphous Co2Ge/CoZr and atomically disordered crystalline GdAl2/CoZr prepared by ball milling take a spin-glass state at low temperatures.[1820] For the ferromagnetic Ni-Mn-Ga alloys, it is found that ball milling not only suppresses the Curie temperature but also weakens the magnetization of the alloy.[12,23] Moreover, the ferromagnetism of the ball-milled Ni-Mn-Ga particles can be restored after annealing at elevated temperatures.[12,23]

Here, the high-energy ball-milling method was used to synthesize Ni50Mn36.7In13.3 fine-sized particles. It was found that the Ni50Mn36.7In13.3 particles transformed to an amorphous-like phase after ball milling for 8 hours. Crystallization of the ball-milled amorphous-like particles during annealing was examined by the in-situ high-energy X-ray total scattering experiments and the differential scanning calorimetry (DSC) techniques. The characterizations of the ball-milled and annealed particles were further studied through the magnetization measurements.

2 Experiment Details

2.1 Sample Preparation

An ingot with the nominal composition of Ni50Mn36.7In13.3 (at. pct) was prepared by arc-melting high-purity elements under an argon atmosphere. The cast ingot was homogenized at 1173 K (900 °C) for 48 hours in vacuum and then quenched in ice water. Some small pieces were cut from the annealed ingot for determining martensitic transformation temperatures. The remainder of the annealed ingot was crushed and ground to powders of ~200 μm in size. Some powders were then annealed at 773 K (500 °C) for 1 hour in vacuum for the in-situ high-energy X-ray diffraction (XRD) experiments. Some powders were further high-energy ball milled using a SPEX-8000 laboratory miller (CertiPrep, Edison, NJ) under an argon atmosphere for 8 hours. The hardened chrome steel balls were used with a ball-to-powder weight ratio of 4:1. During milling, cycles of 10 and 6 minutes rest were performed for avoiding overheating. The total impurity concentration of the ball-milled particles was less than 0.3 at. pct. The ball-milled particles were further studied using in-situ synchrotron X-ray scattering, DSC, a vibrating sample magnetometer (VSM), and superconducting quantum interference magnetometer (SQUID) instruments.

2.2 In-Situ X-Ray Scattering Experiments

The in-situ synchrotron-based high-energy X-ray scattering experiments (115 keV X-rays) were carried out at the 11-ID-C beamline of the Advanced Photon Source, Argonne National Laboratory. During experiments, the X-ray scattering data were collected by a two-dimensional (2-D) image plate detector (Mar345). The CeO2 standard powders were used to calibrate the sample-to-detector distance.

The in-situ high-energy X-ray diffraction experiments were performed for (1) the Ni50Mn36.7In13.3 powders before ball milling and (2) the annealed particles after ball milling, with a sample-to-detector distance of ~1294 mm for attaining a high resolution. The sample was placed into kapton capillaries, and the structural evolution as a function of temperature ranging from 80 K to 500 K (–193 °C to 227 °C) was studied using an Oxford Cryostream system installed at the beamline.

The characterization of the ball-milled particles was carried out, in part, by the in-situ high-energy X-ray total scattering experiments coupled with atomic pair distribution function (PDF) analysis. The 2-D detector was moved to a closer distance from the sample (~246 mm), for collecting scattering data with the maximum wave factor Qmax up to 25 Å−1. The higher value of Qmax is important for the precise PDF analysis.[22] The ball-milled particles were sealed in a Honeywell sample chamber filled with nitrogen installed at the beamline, and the temperature was varied from room temperature (RT) to 684 K (411 °C) step by step. The scattering data were collected at temperatures in the sequence of RT, 373 K, 423 K, 473 K, 523 K, 573 K, 623 K, 673 K, 684 K, 573 K, 373 K (100 °C, 150 °C, 200 °C, 250 °C, 300 °C, 350 °C, 400 °C, 411 °C, 300 °C and 100 °C), and RT, with heating and cooling rates of 10 K min−1 (10 °C min−1). The sample was held for 30 minutes at each temperature, during which five 2-D images were collected for both sample and background. The 2-D X-ray scattering patterns were reduced to one-dimensional (1-D) scattering patterns using the program Fit2D.[24,25] Then, the corresponding real-space atomic PDF was constructed from the 1-D scattering data for analyzing the detailed information on atomic structure.

The atomic PDF method deals with the total scattering, including both Bragg and diffuse components.[22,2629] Thus, it contains the detailed information on the short-, medium-, and long-range atomic configurations. The atomic PDF G(r) is defined as
$$ G\left( r \right) = 4\pi r\left[ {\rho \left( r \right) - \rho_{0} } \right] $$
(1)
where ρ(r) and ρ0 are the local and average atomic number densities, respectively; and r is the radial distance. It peaks at characteristic distances separating pairs of atoms and, hence, reflects the atomic-scale structure.[22] The experimental atomic pair distribution function, G(r), is obtained by a Fourier transformation of the total structure function S(Q) determined through the aforementioned scattering experiments; i.e.,
$$ G\left( r \right) = \frac{2}{\pi }\int_{0}^{\infty } {Q\left[ {S\left( Q \right) - 1} \right]} \sin QrdQ $$
(2)
where Q = 4πsinθ/λ is the magnitude of the wave factor. In this article, the total structure function S(Q) and PDF G(r) were calculated using PDFgetX2,[30] where standard corrections were applied according to the image plate geometry.[31] The refinement of the structure models to the PDF data was performed with the program PDFgui.[32]

2.3 DSC, VSM, and SQUID Measurements

The phase transformation properties of the alloy before and after ball milling were measured in a DSC apparatus (TA instrument, Q100), with heating and cooling rates of 10 K min−1 (10 °C min−1). The RT magnetization curves of the ball-milled and annealed particles were measured by a VSM under different magnetic fields applied up to 2 T. The low-field (0.01 T) thermomagnetization properties of the ball-milled and annealed particles were examined by using a SQUID in the temperature interval of 25 K to 320 K (–248 °C to 47 °C), with the temperature heating and cooling rate of 3 K min−1 (3 °C min−1).

3 Results and Discussion

3.1 Martensitic Transformations and Crystal Structures Before Ball Milling

Figure 1 presents the DSC traces for the Ni50Mn36.7In13.3 as-cast alloy before ball milling. The martensitic transformation start temperature (Ms), the martensitic transformation finish temperature (Mf), the austenitic transformation start temperature (As), and the austenitic transformation finish temperature (Af) were determined to be 373 K, 363 K, 375 K, and 386 K (100 °C, 90 °C, 102 °C, and 113 °C), respectively. The corresponding phase transformation interval (differences between Ms and Mf or As and Af) and thermal hysteresis (difference between Af and Ms) are ~10 K and ~13 K (10 °C and 13 °C), respectively. The enthalpy change (ΔH) around the martensitic transitions calculated from the calorimetry data is ~18.73 J g−1.
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Fig. 1

Constant-rate (10 K min−1 [10 °C min−1]) DSC traces for the Ni50Mn36.7In13.3 alloy before ball milling

According to the transformation temperatures, the Ni50Mn36.7In13.3 original alloy should be in the martensite state at RT. Crystal structures of the martensite and parent phases were studied by the in-situ high-energy X-ray diffraction experiments. Figure 2 gives the 1-D diffraction patterns of the alloy before ball milling, collected at 400 K and 290 K (127 °C and 17 °C) during cooling. At 400 K (127 °C), the alloy exhibits a parent Heusler L21 face-centered-cubic (fcc, space group: Fm\( \overline{3} \)m) structure with lattice parameters of ap = 6.009 Å, as seen from the superstructure reflections of the L21 phase. When cooling to 290 K (17 °C), the alloy presents a modulated martensite structure. Here, the martensite can be indexed as a monoclinic 14 M (seven-layered) modulated structure with the lattice parameters of a14M = 4.412 Å, b14M = 5.602 Å, c14M = 30.401 Å, and β14M = 93.38 deg. This is consistent with that reported by Krenke et al.[33] Such a martensite structure is also observed in Ni-Mn-Ga and Ni-Mn-Sn alloys.[34,35]
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Fig. 2

(Color online) High-energy XRD patterns for the Ni50Mn36.7In13.3 alloy before ball milling, collected at temperatures of 400 K (127 °C) (lower) and 290 K (17 °C) (upper) during cooling. The indexing of Bragg reflections corresponds to the Heusler parent and 14 M martensite, respectively. The lower-inset refers to a Heusler structure with Ni, Mn, and Ga atoms at respective sites colored in blue, orange, and pink

3.2 Structural Evolution of the Ball-Milled Particles During Annealing

Structure transitions of the ball-milled Ni50Mn36.7In13.3 particles during annealing were studied by the in-situ high-energy X-ray total scattering experiments. Figure 3 shows the high-energy X-ray scattering patterns for the ball-milled particles collected at different annealing temperatures. As discussed in Section III–A, at RT, the Ni50Mn36.7In13.3 original alloy shows a 14 M modulated martensite structure, exhibiting sharp well-defined Bragg peaks in the diffraction pattern. For the particles ball milled for 8 hours, there are only two broad diffraction peaks in the scattering patterns of the particles collected at RT (Figure 3), taken as evidence for a main “X-ray amorphous” phase.[21] The amorphous nature of the ball-milled particles was confirmed by the subsequent PDF analysis and DSC measurements. The ball-milling process resulted in a variety of defect structures (dislocations, vacancies, stacking faults, grain boundaries, etc.). These defect structures can destabilize the ordered nature of the lattice and raise the free energy of the system, leading to the formation of an amorphous phase.[17] The somewhat triangular shape of the first broad peak is probably due to a disordered solid solution reflection superimposed on the reflection of the amorphous phase.[21] As the Ni-Mn-In FSMAs undergo an order (L21) to disorder (B2) transition during heating before they melt, milling of this kind of alloys should first result in the formation of a disordered solid solution (with the crystalline structure) and then the amorphous phase.[17] Thus, it is expected that after ball milling for 8 hours, the main structure of the particles is amorphous, but mixing with disordered crystalline structure. This is defined as the amorphous-like phase here.
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Fig. 3

High-energy X-ray scattering patterns for the Ni50Mn36.7In13.3 ball-milled particles collected at a series of temperatures during in-situ annealing, showing structural transitions from an amorphous-like structure (RT) first to a B2 structure (523 K [250 °C] for 15 min) and then to a Heusler structure (684 K [411 °C] for 30 min). The tiny peaks marked by open circles may come from a different phase

During heating, the amorphous-like phase remains stable up to 473 K (200 °C). When the temperature increases to 523 K (250 °C), the broad shape of the diffraction peaks remains stable for 2 to 3 minutes. During further annealing at this temperature, the diffraction peaks become sharp and more well-defined peaks appear. After annealing at 523 K (250 °C) for 15 minutes, the Bragg diffraction peaks in the scattering pattern, as shown in Figure 3, can be indexed as a disordered body-centered-cubic (bcc) B2 structure rather than the ordered Heusler L21 structure. The B2 structure of the particles after annealing at 523 K (250 °C) can be confirmed by the subsequent atomic PDF structure refinement. This B2 structure remains stable up to 623 K (350 °C). With further annealing at 673 K and 684 K (400 °C and 411 °C), some small peaks appear between some main peaks belonging to B2. The collected scattering pattern for the particles after annealing at 684 K (411 °C) for 30 minutes is shown in Figure 3. Most of the Bragg peaks can be indexed by the Heusler L21 structure, except for a few tiny peaks (indicated by open circles), which may result from the eutectic separation of a second phase. Further high-resolution XRD patterns (discussed in Section III–C) proved that this Heusler structure is a slightly tetragonal distorted structure.

As the ball-milled particles are of limited structural coherence, accurate and detailed information on structures was further extracted through the corresponding real-space PDF analysis. Figures 4(a) through (d) give the calculated experimental atomic PDFs for (a) the CeO2 standard powders, (b) ball-milled particles at RT, (c) ball-milled particles after annealing at 523 K (250 °C) for 15 minutes, and (d) ball-milled particles after annealing at 684 K (411 °C) for 30 minutes. The atomic PDF for the CeO2 standard powders extended to a real-space distance as high as ~50 Å displays a well-defined crystalline structural feature (Figure 4(a)), determined by the resolution of used X-ray scattering instrument. The experimental PDF for the ball-milled Ni-Mn-In particles at RT (Figure 4(b)) shows the typical characterization for amorphous metals and alloys.[26,36] This confirms the main amorphous nature of the ball-milled particles. The experimental PDFs for the particles after annealing at 523 K (250 °C) become rich in well-defined crystalline structural features (Figure 4(c)), but they vanish at a shorter interatomic distance of ~40 Å compared with that for CeO2. This is due to the limited structural coherence of the particles, which is related with the grain size or atomic disorder effects resulting from the ball-milling process.[17] After further annealing at 684 K (411 °C) for 30 minutes, the particles present almost the same experimental PDF (Figure 4(d)) as that for the particles after 523 K (250 °C) annealing. Only a small shift of positions and slight change of intensities in some peaks can be found in the PDFs for the particles annealed at two different temperatures.
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Fig. 4

Experimental atomic PDFs for (a) CeO2 standard powders and Ni50Mn36.7In13.3 ball-milled particles during in-situ annealing: (b) at RT, (c) after annealing at 523 K (250 °C) for 15 min, and (d) after annealing at 684 K (411 °C) for 30 min. The PDFs are extracted from the corresponding scattering patterns shown in Fig. 3. The decay in the PDF intensity for CeO2 is due to the finite resolution of the high-energy X-ray

The PDF refinement of atomic structures was conducted for the 523 K (250 °C) annealed particles using the PDFGui code. The PDFGui employs a least-squares procedure to compare experimental and model data (PDF) calculated from a plausible structural model.[32] Figure 5 gives the refinement results, based on the B2 (disordered bcc) structural model. It can be seen that the experimental PDF for the 523 K (250 °C) annealed particles (open circles) can be well reproduced by the B2-modeled PDF (solid line) over a wide radial distance up to 25 Å. The resulting reliability factor (Rw) is ~12.9 pct. As a consequence, the crystal structure of the 523 K (250 °C) annealed particles can be proved to be B2 (space group: Pm\( \overline{3} \)m), with Ni occupying the center of the cube and Mn/In distributed randomly at the cube corners. The corresponding refined lattice parameter is aB2 = 3.009 Å.
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Fig. 5

(Color online) Experimental (symbols) and model (lines) atomic PDFs G(r) for the Ni50Mn36.7In13.3 ball-milled particles after annealing at 523 K (250 °C), with the difference given in the lower part of the plots. The model PDF is based on the cubic B2 structure (refer to the inset). The corresponding reliability factor (Rw) is ~12.9 pct

Figure 6 presents a comparison between the PDFs for the ball-milled particles (1) at RT (dashed line) and (2) after annealing at 523 K (250 °C) (solid line), over the first few atomic shells in the low-r region (r < 6 Å). According to the preceding discussion, the ball-milled particles were amorphous at RT and transformed to the B2 crystalline structure after annealing at 523 K (250 °C). By using the Gauss least-squares fitting method, the peak positions can be obtained, which provide information on the interatomic distances.[2629] The intensity of the PDF G(r) reflects the statistical number of atom pairs at a given distance.[2629] For the B2 structure (solid line), the first peak in the PDF is positioned at approximately 2.61 Å (~\( \sqrt{3/2} \)·aB2), which is the first neighbor Ni-Mn and Ni-In atomic distance. The second PDF peak located at ~2.98 Å (~aB2) corresponds to the first neighbor Mn-In, Mn-Mn, and In-In atomic distance. The PDFs for the amorphous and B2 structure of the particles are quite similar in the r region near the first peak but differ in positions and intensities in the extended r region. This similarity in PDFs near the first peak suggests that the first atomic shell is the same for the amorphous and the B2 structures of the particles.
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Fig. 6

Comparison between the PDFs for the ball-milled particles at RT (dashed line) and after annealing at 523 K (250 °C) (solid line), in the low-r region (r < 6 Å)

Figure 7 shows the constant-rate (10 K min−1 [10 °C min−1]) DSC traces for the ball-milled amorphous particles, measured in the temperature range from 273 K to 673 K (0 °C to 400 °C) for two cycles. During heating in the first cycle (solid line), a huge and sharp exothermic peak appears at ~523 K (250 °C), which should correspond to the amorphous to B2 crystallization process. The calculated crystallization enthalpy is ~63.5 J g−1. This is additional evidence for the amorphous-like nature of the ball-milled particles. The glass transition is not observed prior to crystallization as that in many ball-milled amorphous systems.[17,21] This may be due to the fact that there exist a lot of defect structures in the ball-milled particles, which would induce some exothermic or endothermic reactions during heating prior to crystallization and, consequently, obscure the glass transition. During the subsequent cooling in the first cycle, an exothermic peak appears at around 313 K (40 °C). This should correspond to the martensitic transformation, as the particles were crystallized to B2 structure during heating in the first cycle. In the DSC curve of the second cycle (dashed line), both the martensitic and austenitic transformation peaks can be observed. The transformation temperatures Ms, Mf, As, and Af are determined to be 320 K, 289 K, 310 K, and 340 K (47 °C, 16 °C, 37 °C, and 67 °C), respectively. The corresponding phase transformation interval and thermal hysteresis are ~30 K and ~23 K (30 °C and 23 °C), respectively. Both are larger than that for the alloy before ball milling. The enthalpy change (ΔH) around the martensitic transitions calculated from the calorimetry data is ~5.80 J g−1, about 1/3 of the value for the alloy before ball milling.
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Fig. 7

(Color online) Constant-rate (10 K min−1 [10 °C min−1]) DSC traces for the Ni50Mn36.7In13.3 ball-milled particles, measured in the temperature range from 273 K to 673 K (0 °C to 400 °C) for two cycles

3.3 Structural Transitions of the Annealed Particles Under Low Temperature

As discussed in Section III–B, the Ni50Mn36.7In13.3 amorphous particles were obtained after ball milling for 8 hours. In-situ annealing of the ball-milled amorphous particles first led to their crystallization with the formation of a B2 structure after annealing at 523 K (250 °C) for 15 minutes. After further annealing at 684 K (411 °C) for 30 minutes, most of the Bragg peaks in the scattering pattern can be indexed by the Heusler L21 structure (Figure 3). During subsequent in-situ cooling from 684 K (411 °C) to RT, the scattering patterns remained the same at temperatures of 684 K, 573 K, and 373 K (411 °C, 300 °C, and 100 °C). There appeared to be some changes in the scattering pattern when cooled to RT, which was supposed to be the occurrence of a martensitic transition from the Heusler-type parent phase to martensite. On the other hand, the DSC curves shown in Figure 7 indicated that a martensitic transition from the B2-type parent phase to martensite happened at the temperature range of 320 K to 289 K (47 °C to 16 °C). In this section, low-temperature structural transitions of the annealed particles with both B2-type and Heusler-type parent phase are studied as a function of temperature from 400 K to 120 K (127 °C to –153 °C). The B2-type particles are obtained by in-situ annealing the ball-milled particles at 523 K (250 °C) for 30 minutes, whereas the Heusler-type particles are obtained by in-situ annealing the ball-milled particles at 684 K (411 °C) for 30 minutes.

Figure 8(a) shows the high-energy XRD patterns for the 523 K (250 °C) annealed particles collected at various temperatures during cooling from 400 K to 120 K (127 °C to –153 °C). At higher temperatures, the 523 K (250 °C) annealed particles show the B2 structure for the parent phase. During cooling, a martensite phase with a modulated 10 M (five-layered) structure appears at ~320 K (47 °C) and becomes stable below 250 K (–23 °C). This martensite phase has a monoclinic structure (indexing in Figure 8(a)), with the lattice parameters of a10M = 4.359 Å, b10M = 5.747 Å, c10M = 21.648 Å, and β = 90.33 deg. A similar monoclinic 10 M modulated structure of martensite has also been found in Mn-rich Ni-Mn-Ga alloys.[37] During heating, the reverse martensitic transformation from the monoclinic 10 M martensite to B2 parent phase begins at ~270 K (–3 °C) and finishes at ~340 K (67 °C).
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Fig. 8

High-energy XRD patterns at various temperatures for in-situ structural characterization of phase transformation behaviors in Ni50Mn36.7In13.3 particles: (a) ball-milled and 523 K (250 °C) annealed particles and (b) ball-milled and 684 K (411 °C) annealed particles. The indexing of Bragg reflections in (a) and (b) corresponds to 10 M martensite structure and slightly tetragonal distorted Heusler structure, respectively. The tiny peaks marked by open circles in (b) may come from a different phase

Figure 8(b) shows the high-energy XRD patterns for the 684 K (411 °C) annealed particles collected at various temperatures during cooling from 400 K to 120 K (127 °C to –153 °C). The XRD patterns remain the same during cooling from 400 K to 330 K (127 °C to 57 °C). According to the high-energy X-ray total scattering patterns and the aforementioned discussions, the 684 K (411 °C) annealed particles exhibit a Heusler-type structure at 400 K (127 °C). The high-energy XRD pattern collected at 330 K (57 °C) (or 400 K [127 °C]) is shown in Figure 8(b). Different from the diffraction patterns observed for the stoichiometric Heusler structure, most of the main peaks split into two different peaks, as seen from the high-resolution XRD. This can be explained by considering a small tetragonal distortion of the Heusler structure, with the lattice parameters of a = 6.09 Å (slightly elongating 1.16 pct) and c = 5.99 Å (slightly contracting 0.50 pct). It should be noted that the splitting of the peaks was not visible in Figure 3, which can be attributed to the low resolution of the total scattering experiments. The sample-to-detector distance is much closer during the total scattering experiments, which gives rise to a higher Qmax value but at the same time lowers the reciprocal-space resolution. The indexing based on the tetragonal distorted Heusler structure is shown in Figure 7(b), with a few tiny peaks (marked by open circles) reflecting the second phase. During further cooling from 330 K (57 °C), the diffraction patterns begin to change at ~320 K (47 °C) and become stable below 260 K (–13 °C). The XRD patterns collected at 260 K and 120 K (−13 °C and –153 °C) are shown in Figure 8(b). Apparently, there is a structural transition during cooling from 330 K to 260 K (57 °C to −13 °C), as seen from the changes of the relative peak intensities and the appearances of some new peaks (marked by arrows). The nature of this low-temperature phase needs further study. When heating, the reverse structural transition (from the low-temperature phase to the tetragonal distorted Heusler) begins at ~320 K (47 °C) and finishes at ~350 K (77 °C).

3.4 Magnetic Properties of the Ball-Milled and Annealed Particles

Figure 9(a) shows the RT magnetization curves obtained by the VSM technique for the Ni50Mn36.7In13.3 alloy before ball milling (solid squares), ball-milled amorphous particles (open circles), 523 K (250 °C) annealed particles (open triangles), and 684 K (411 °C) annealed particles (solid diamonds). The curve for the alloy before ball milling exhibits a typical nonmagnetic behavior, with the magnetization close to zero even at a magnetic field strength of 2 T. As indicated by Sutou et al.[38] and Ito et al.,[39] the original bulk alloy with a Ni50Mn36.7In13.3 atomic composition should be para- or antiferromagnetic, whether the alloy is at the martensite or parent state. Certainly, here the Ni50Mn36.7In13.3 alloy before ball milling shows a modulated martensite structure at RT (Section III–A). The curve for the ball-milled amorphous particles also exhibits a nonmagnetic behavior at RT, showing a similar magnetization to the alloy before ball milling. Annealing of the amorphous particles at 523 K (250 °C) for 15 minutes led to their crystallization with a B2 structure. The RT magnetization for the 523 K (250 °C) annealed particles becomes larger, compared with that for the ball-milled amorphous particles. However, the curve still shows a para- or antiferromagnetic-like behavior, as shown in Figure 8(a). Further annealing of the particles at 684 K (411 °C) for 30 minutes led to the formation of a Heusler structure with a small tetragonal distortion. The RT magnetization curve for the 684 K (411 °C) annealed particles exhibits a ferromagnetic-like (or ferrimagnetic-like) behavior at RT. Here, the increased magnetization is consistent with the increased atomic ordering.
https://static-content.springer.com/image/art%3A10.1007%2Fs11661-011-0724-x/MediaObjects/11661_2011_724_Fig9_HTML.gif
Fig. 9

(Color online) (a) Room-temperature magnetization curves for the Ni50Mn36.7In13.3 alloy before ball milling (solid squares), ball-milled and unannealed particles (open circles), ball-milled and 523 K (250 °C) annealed particles (open triangles), and ball-milled and 684 K (411 °C) annealed particles (solid diamonds). The temperature dependences of magnetization curves for the ball-milled and differently annealed particles are shown in (b) through (d), respectively

The thermomagnetization curves obtained with SQUID for the ball-milled amorphous particles, 523 K (250 °C) annealed particles, and 684 K (411 °C) annealed particles are given in Figures 9(b) through (d), respectively. Both the zero-field-cooled (ZFC) and field-cooled (FC) properties are measured at a magnetic field strength of H = 0.01 T, with the temperature interval of 25 K ≤ T ≤ 320 K (−248 °C ≤ T ≤ 47 °C). There exists a common feature for the thermomagnetization curves shown in Figures 9(b) through (d); that is, the ZFC curves behave differently with the FC curves at low temperatures. As illustrated in Figure 9(b) for the ball-milled amorphous particles, with decreasing temperature, the FC cooling and heating magnetizations increase gradually until to the lowest measuring temperature, whereas the ZFC heating magnetization first increases gradually with decreasing temperature and decreases suddenly after showing a cusp at ~60 K (–213 °C). Such a spitting behavior of FC and ZFC curves is typically observed in the spin-glass systems,[1820] with the cusp representing the freezing temperature Tf. The thermomagnetization curves for the 523 K (250 °C) annealed particles also show a typical spin-glass behavior with Tf ~ 140 K (–133 °C), as shown in Figure 9(c). According to the aforementioned discussions, the 523 K (250 °C) annealed particles present an atomically disordered B2 structure at higher temperatures and transform to the 10 M modulated martensite structure with decreasing temperature from 320 K to 250 K (47 °C to –23 °C). Here, it should be noted that both B2 parent phase and 10 M martensite are paramagnetic-like, as shown in Figure 9(c). Both ball-milled amorphous particles and 523 K (250 °C) annealed crystalline particles with the disordered B2 structure exhibit a transition from the paramagnetic-like to the spin-glass state with decreasing temperature. The freezing temperature of the 523 K (250 °C) annealed particles (140 K [–133 °C]) is higher than that of the ball-milled amorphous particles (60 K [−213 °C]). Zhou and Bakker[20] reported that both ball-milled amorphous and atomically disordered crystalline CoZr take a spin-glass state at low temperatures. They found that the Tf value of the disordered crystalline CoZr is higher than that of the amorphous CoZr, which was attributed to the larger magnetic cluster sizes in the disordered crystalline CoZr.

For the 684 K (411 °C) annealed particles (Figure 9(d)), as the temperature decreases, the ZFC and FC magnetizations drastically increase between 320 K and 300 K (47 °C and 27 °C) owing to the transformation from paramagnetic to ferromagnetic (or ferrimagnetic) with a Curie temperature Tc ≈ 310 K (37 °C). As the temperature further decreases from 300 K (27 °C), the FC cooling and heating magnetizations still increase but with a decreasing slope. On the other hand, the ZFC heating magnetization begins to decrease slowly from 300 K (27 °C) down to the lowest measuring temperature. Compared with Figures 9(b) and (c), the splitting of the ZFC and FC curves for the 684 K (411 °C) annealed particles (Figure 9(d)) is not typical for a spin-glass behavior, which is expected to be associated with coexisting ferromagnetic and antiferromagnetic exchanges. This feature can be attributed to an increase in the ferromagnetic (or ferrimagnetic) component for the particles after annealing at 684 K (411 °C). Such a separation between ZFC and FC curves has also been found in the deposited Ni2MnIn alloy films.[40] In addition, as discussed in Section III–C, the 684 K (411 °C) annealed particles show a tetragonal distorted Heusler structure at higher temperatures. During cooling, the structural transition from the tetragonal distorted Heusler to a low-temperature phase happens in the temperature range of 320 K to 260 K (47 °C to –13 °C); during heating, the reverse structural transition happens in the temperature range of 320 K to 350 K (47 °C to 77 °C). In combination with the thermomagnetization measurements shown in Figure 9(d), it can be concluded that the tetragonal distorted Heusler phase is paramagnetic, whereas the transformed low-temperature phase is ferromagnetic (or ferrimagnetic) with the Curie temperature Tc ≈ 310 K (37 °C) close to the structural transition temperatures.

4 Conclusions

In conclusion, the Ni50Mn36.7In13.3 amorphous-like particles were synthesized through the high-energy ball-milling method. Crystallization of the amorphous particles during annealing was studied by the in-situ high-energy X-ray total scattering experiments and PDF analysis. During annealing, the particles first transform from amorphous to an atomically disordered B2 crystalline structure, and then to an ordered Heusler L21 structure with a small tetragonal distortion. The annealed crystalline particles undergo different structural transitions during cooling, tailored by the atomic order. Magnetic measurements show that both amorphous and atomically disordered crystalline particles take a spin-glass state at low temperatures, whereas the ordered crystalline particles exhibit a ferromagnetic behavior with the Curie temperature Tc ≈ 310 K (37 °C).

Acknowledgments

This work is supported by the National Natural Science Foundation of China (Grant Nos. 50725102 and 50971031). Use of the Advanced Photon Source was supported by the United States Department of Energy, Office of Science, Office of Basic Energy Science, under Contract No. DE-AC02-06CH11357.

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© The Minerals, Metals & Materials Society and ASM International 2011