Microstructures and Magnetic Properties of Fe70−xPd30Nix High-Temperature Ferromagnetic Shape Memory Alloys

1 23

Journal of Superconductivity and Novel Magnetism

Incorporating Novel Magnetism ISSN 1557-1939

J Supercond Nov Magn

DOI 10.1007/s10948-014-2690-1

Microstructures and Magnetic Properties

of Fe

_70−x

Pd

₃₀

Ni

_x

High-Temperature

Ferromagnetic Shape Memory Alloys

Yin-Chih Lin

1 23

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J Supercond Nov Magn DOI 10.1007/s10948-014-2690-1

ORIGINAL PAPER

Microstructures and Magnetic Properties of Fe

70

−x

Pd

30

Ni

x

High-Temperature Ferromagnetic Shape Memory Alloys

Yin-Chih Lin

Received: 13 June 2014 / Accepted: 8 August 2014 © Springer Science+Business Media New York 2014

Abstract This research paper mainly presents an

investi-gation of the microstructures and magnetic properties of bulk ferromagnetic shape memory (FSM) Fe70−xPd30Nix

(Nix = 4, 8 at.%) alloys, by transmission electron

microscopy (TEM), a magnetostrictive-meter setup, and a superconducting quantum interference device (SQUID) magnetometer. The FSM alloys were homogenized through hot and cold strain forging (SF) to a∼38 % reduction in thickness, solution-treated (ST) with annealing recrystal-lization at 1100◦C or 8 h, quenched in ice brine, and then aged at 500◦C for 100 h (5 HTA). The investigation of the microstructures and magnetic properties indicated that the higher Ni content (Nx = 8 at.%) in the Fe62Pd30Ni8alloy

SF and ST reduced the saturation magnetostriction at RT. However, with higher Ni content in the Fe62Pd30Ni8alloy,

the decomposition of L10+ L1mtwin phases into

stoichio-metric L10 + L1m+ αbct structures was suppressed after

the alloy was ST and aged at 5 HTA, as confirmed by TEM investigation. The result was that the FSM Fe62Pd30Ni8

alloy maintained a high saturation magnetostriction and magnetization after the alloy was ST and aged at 5 HTA. This magnetic property of the Fe62Pd30Ni8 FSM

alloy makes it suitable for application in high-temperature (T < 500 ◦C) and high-frequency environments. How-ever, low Ni content FSM Fe66Pd30Ni4 (Nx = 4 at.%)

alloy SF, ST, and aged at 5 HTA underwent decomposi-tion of the L10+ L1m twin phases into the stoichiometric

L10+ L1m+ αbctstructures, as confirmed by TEM, leading

to a decrease of saturation magnetostriction and magnetiza-tion. This magnetic property of the Fe66Pd30Ni4FSM alloy

Y.-C. Lin ()

Department of Mold and Die Engineering, National Kaohsiung University of Applied Sciences, 807-78, Kaohsiung, Taiwan e-mail: [email protected]

is not suitable for application in high-temperature (T < 500◦C) environments.

Keywords FSM Fe70−xPd30Nix(Nix= 4, 8 at.%) alloys ·

Aging treatment· TEM microstructure · Magnetic property

1 Introduction

Magnetic field control of high-temperature shape memory effects was recently suggested as a principle for the opera-tion of a new type of actuator material. At present, a new class of high-temperature shape memory alloys (HTSMAs) has been widely investigated due to the ability of the mate-rials to suppress transitions at high temperatures (exceeding 400◦C). This magnetic property makes HTSMAs suitable for actuation in high-temperature environments commonly found in the aerospace, automotive, and oil industries or in certain other devices [1–3]. This study investigated in detail the effects of varying the amount of Ni in Fe70−xPd30Nix

(Nix = 4, 8 at.%) alloys in samples strain-forged (SF) and

solution-treated (ST) before being aged at 5 HTA. The influ-ences of high-temperature (5 HTA) aging treatment on the transmission electron microscopy (TEM) microstructures, phase transition behavior, saturation magnetostriction, and

M-H curves of the materials were investigated in detail in order to evaluate their thermal stability for high-temperature applications.

2 Experimental Procedure

Fe66Pd30Ni4 and Fe62Pd30Ni8 (at.%) ferromagnetic shape

memory alloys (FSMAs) were prepared by melting pure electrolytic iron (99.9 %), pure palladium (99.95 %), and

J Supercond Nov Magn

pure nickel powder (99.95 %) in an arc vacuum furnace under a controlled protective argon atmosphere. Samples were sliced from the cast ingot and sealed in an evacuated quartz capsule, where they were homogenized at 1050◦C for 70 h, followed by quenching in ice brine. After homog-enization, the specimens were SF to a ∼38 % reduction in thickness, polished, sealed in an evacuated quartz cap-sule again, and then ST and annealed for recrystallization at 1100◦C for 8 h, followed by quenching in ice brine. After that, parts of the samples were aged at 5 HTA. Thin foils for TEM studies were prepared by double-jet electropolishing in a solution containing 75 % acetic acid, 15 % perchloric acid, and 10 % methanol in a temperature range of−7◦C∼ 10◦C using a current density of 2 to 4 A/cm2. Transmission electron microscopy, with a double-tilt stage, was performed in an analytical-type high-resolution electron microscope (JEM-2100F TEM) with a field emission gun operating at 200 kV. The magnetostriction measurements were taken with a strain gauge and Vishay Micro-Measurements Model P3 strain indicator and recorder. A sample with dimensions of 15 mm (length)× 5 mm (width) × 1 mm (thickness) was used for the magnetostriction measurement. A FLA-2-11-A515411-type strain gauge from TML (Tokyo Sokki Kenkyujo Co. Ltd.) was employed. The gauge was glued to the sample along its longitude with the M-Bond 200 Adhe-sive Kit. To cure the glue, both the gauge and the specimen were kept at RT for 2 days, after which the magnetostrictive strains were measured. The hysteresis loop measurements were performed with a superconducting quantum interfer-ence device (SQUID) magnetometer. The magnetization measurements were performed on a specimen with dimen-sions of 5 mm (length)× 1 mm (width) × 1 mm (thickness). The magnetization versus magnetic field (M− H ) curves for the samples were measured at room temperature (300 K) with a maximum applied field of 10,000 Oe. Solution treatment means that the alloy sample which may be in the wrought or cast form is heated to a temperature between the solvus and solidus temperatures and soaked there until a uniform solid-solution structure is produced. Then, the sam-ple is rapidly cooled to a lower temperature, and the cooling medium is usually water or ice brine at room temperature. The structure of the alloy sample after ice brine quenching consists of a supersaturated solid solution.

3 Results and Discussion

3.1 TEM Microstructures of the SF Fe66Pd30Ni4

and Fe62Pd30Ni8Alloys ST and the SF Alloys Post ST

then Aged at 5 HTA (500◦C/100 h)

The essential TEM images of the Fe66Pd30Ni4 FSM alloy

SF to a∼38 % reduction in thickness, annealed at 1100◦C

for 8 h, and quenched in ice brine are shown in Fig.1a–f. Shown in Fig. 1a is the selected area diffraction pat-tern (SADP) of zone axis [0 1 1]L10//[4 1 1]L1m

(hkl denotes tetragonal L10 structure and the ordered L10

martensitic structure with lattice parameters of a= 3.807 ˚A,

c= 3.657 ˚A, and c/a = 0.96; hkl denotes L1mmonoclinic

phase with lattice parameters of a = 3.147 ˚A, b = 3.807 ˚

A, c = 3.113 ˚A, and β = 92.318◦). The dark field (DF) micrographs of Fig.1b were obtained with the diffraction vector g = [200]L10 corresponding to Fig.1a. In this DF

image, the tetragonal g(200)L10 twinning plates seem to

be aligned alternately and form a lamellar twin structure. Figure1c is a DF image using the (1 2 2)L1mreflection

corresponding to Fig. 1a. It is found that the L1m

twins (bright contrast) were irregular, and this twinning mode is identified as{0 1 1}L1m. The L10+ L1m twin

phases had better saturation magnetostriction, as shown in Fig. 5a. Figure 1d is a bright field (BF) image. Shown in Fig. 1e is a nano beam diffraction pattern (NBDP) of zone axis[ 1 1 1]L10//[ 1 0 1]L1mtaken from the

same sample. For confirmation that tetragonal L10 and

monoclinic L1mphase were present in the specimen, a

high-resolution TEM (HRTEM) image, taken from Fig. 1d, is shown in Fig.1f. Careful measurement of the lattice space revealed that the d spacing of the tetragonal L10 structure

was 0.26 nm and the d spacing of the monoclinic L1mphase

was 0.22 nm; therefore, the plane (Fig.1f) can be reasonably inferred to be (101)L10, and (101)L1m, respectively. The

tetragonal L10 and monoclinic L1m structures existed in

the Fe66Pd30Ni4FSM alloys, as confirmed consistently by

NBDP and HRTEM [4–6].

Fig. 1 TEM images of the Fe66Pd30Ni4 (at.%) alloy strain-forged

(SF) to a ∼38 % reduction in thickness, ST, and recrystallization-annealed at 1100◦C for 8 h: a SADP of zone axis[0 1 1]L10//

[ 4 1 1]L1m (hkl denotes tetragonal L10 reflection; and hkl

denotes L1mmonoclinic structure), b DF image of g = [200]L10, c

DF image of g = [1 2 2]L1m, d BF image, e NBDP of zone axis

[ 111]L10//[ 101]L1m, and f HRTEM image showing d spacing of

the (101)L10and (101)L1m, respectively

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Figure 2a–f is a series of TEM images taken from the SF Fe66Pd30Ni4 FSM alloy post ST and aged at

5 HTA. Shown in Fig. 2a is the SADP of zone axis [ 1 0 1]L10//[ 4 1 5]L1m//[2 3 3]αbct (hkl

denotes tetragonal L10 reflection; hkl denotes L1m

mon-oclinic phase; and hkl denotes αbct structure with lattice

parameters of a = 2.894 ˚A, c = 2.876 ˚A, c/a = 0.994). The decomposition of the L10+ L1mtwin phases into

stoi-chiometric L10+ L1m+ αbctstructures is clearly shown in

the SADP image. Figure2b is a BF image. A comparison of the BF micrographs in Figs.1d and2b revealed distinct differences in the microstructures of L10+L1mtwin phases

and stoichiometric L10+ L1m+ αbct structures, in which

the latter structures caused a drastic decrease in saturation magnetostriction, as shown in Fig.5c. A higher magnifica-tion BF image of Fig.2b is shown in Fig.2c, in which the stoichiometric L10structure and L1m+αbctphases are

indi-cated by arrows, respectively. Figure2d is an NBDP taken from the area of Fig.2c marked with a circle A, revealing

(001)αbct; (011)L1m; (101)αbct; (2 0 1)L1m, and (002)αbct

reflection. Shown in Fig.2e, f are NBDP taken from the areas of Fig. 2c marked with circles B and C, indicating zone axes of[0 1 1]L10and[301]L10, respectively. The

stoichiometric L10+ L1m+ αbctstructures existed in the 5

HTA aged Fe66Pd30Ni4alloys, as confirmed by the NBDP

of TEM.

Figure 3a–f shows a series of TEM images of Fe62Pd30Ni8 FSM alloy SF to a ∼38 % reduction in

thickness and thermally ST at 1100◦C for 8 h. Figure3a is an SADP with zone axis[ 1 0 1]L10//[0 1 0]L1m

Fig. 2 TEM micrographs of the Fe66Pd30Ni4FSM alloy SF and ST

then aged at 5 HTA: a SADP showing zone axis[ 1 0 1]L10//

[ 4 1 5]L1m//[2 3 3]αbct(hkl denotes tetragonal L10

reflec-tion; hkl denotes L1mmonoclinic phase; and hkl denotes αbct

struc-ture), b BF image, c higher magnification of BF image, d NBDP taken from (c) area marked with a circle A, e NBDP of zone axis

[0 1 1]L10 taken from (c) area marked with a circle B, and f

NBDP of zone axis [301]L10 taken from (c) area marked with a

circle C

(hkl denotes a tetragonal L10structure with lattice

param-eters of a = 3.808 ˚A, c = 3.633 ˚A, and c/a = 0.954; hkl denotes L1m monoclinic phase with lattice parameters

of a = 3.128 ˚A, b = 3.808 ˚A, c = 3.111 ˚A, and β = 92.187◦). Figure3b is a DF image using the (1 1 1)L10

reflection corresponding to Fig. 3a. In this DF image, the bright contrast shows the martensitic L10 twin structures,

which contribute to saturation magnetostriction in the alloys that may be useful in high-temperature actuator application. A DF image using the (3 0 3)L1mreflection

correspond-ing to Fig. 3a is shown in Fig. 3c, in which the bright contrast represents the L1mtwin structures. These

marten-sitic L10 +L1m twins provide a lattice invariant shear for

maintaining an invariant habit plane [7–9]; therefore, they are expected to contribute to the ferromagnetic shape mem-ory effect (FSME) in the alloy for use in high temperature magneto-mechanical applications. Figure3d is a BF image. A comparison of Fig.3b with Fig.1b demonstrates that dif-ferent amounts of Ni content in the Fe70−xPd30Nix alloys

generate distinct differences in the microstructures of the martensitic L10 twin phase. Shown in Fig.3e is an NBDP

of the [5 2 4]L10//[2 3 6]L1m zone axis taken from

the same sample. An HRTEM image taken from Fig.3d is shown in Fig.3f. Careful measurement of the lattice space revealed that the d spacing of the tetragonal L10 structure

was 0.216 nm and the d spacing of the monoclinic L1m

martensitic phase was also 0.216 nm; therefore, the planes (Fig. 3f) can be reasonably inferred to be (111)L10 and (101)L1m, respectively.

Shown in Fig. 4a–f is a series of TEM images taken from the SF Fe62Pd30Ni8 alloy post ST and

aged at 5 HTA. Figure 4a is the SADP of zone axis

Fig. 3 TEM images of the Fe62Pd30Ni8alloy SF to a∼38 %

reduc-tion in thickness, ST, and recrystallizareduc-tion-annealed at 1100◦C for 8 h: a SADP of zone axis[ 1 0 1]L10//[0 1 0]L1m(hkl denotes

tetragonal L10reflection and hkl denotes L1mmonoclinic structure),

b DF image of g= [1 1 1]L10, c DF image of g= [3 0 3]L1m,

d BF image, e NBDP of zone axis[5 2 4]L10//[2 3 6]L1m, and

f HRTEM image showing d spacing of the (111)L10and (101)L1m

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Fig. 4 TEM images of the Fe62Pd30Ni8 FSM alloy SF and ST

then aged at 5 HTA: a SADP showing zone axis[ 10 1]L10//

[1 1 1]L1m(hkl denotes tetragonal L10reflection and hkl denotes

L1mmonoclinic phase), b DF image of g= [020]L10, c DF image of

g= [1 10]L1m, d BF image, e NBDP of zone axis[010]L1mtaken

from area in (f) marked with circle B, and f HRTEM image taken from area in (d) marked with circle A, showing d spacing of the (111)L10

and (1 0 1)L1m, respectively

[ 1 0 1]L10//[1 1 1]L1m. The αbct structure is not

evident in the SADP, indicating that the decomposition of L10+L1mtwin phases into stoichiometric L10+L1m+αbct

structures was suppressed after the 5-HTA aging treatment. The DF image of Fig. 4b was obtained with the diffrac-tion vector g = [020]L10 corresponding to Fig.4a. In this

DF image, the bright contrast is L10 phase. Figure4c is a

DF image using the (1 1 0)L1mreflection corresponding

to Fig.4a, showing the bright contrast of the L1m

mono-clinic structure. Figure4d is a BF image. Shown in Fig.4e

is an NBDP of the[010]L1mzone axis taken from the area

in Fig.4f marked with a circle B. An HRTEM image taken from the area of Fig.4d marked with a circle A is shown in Fig.4f. Careful measurement of the lattice space revealed that the d spacing of the tetragonal L10structure was 0.216

nm and the d spacing of the monoclinic L1m phase was

0.22 nm; therefore, the planes (Fig.4f) can be reasonably inferred to be (111)L10and (1 0 1)L1m, respectively.

3.2 Saturation Magnetostriction of the SF Fe66Pd30Ni4

and Fe62Pd30Ni8FSM Alloys ST and the SF Alloys Post

ST Then Aged at 5 HTA

Figure5a–d shows the saturation magnetostriction (λs)vs.

magnetic field (H ) curves, or λs− H curves, measured at

RT of the (a) Fe66Pd30Ni4and (b) Fe62Pd30Ni8FSM alloys

SF to a∼38 % reduction in thickness, ST at 1100◦for 8 h, and quenched in ice brine. The saturation magnetostric-tion λswas evaluated as λs = (2/3) × [(λ||)− (λ⊥)], where λ_|| (λ⊥)is the magnetostriction in the longitudinal direc-tion with a magnetic field parallel (perpendicular) to the sample’s length. It can be seen in Fig.5a, b that the satura-tion magnetostricsatura-tion of the SF and ST Fe66Pd30Ni4 alloy

(64.7×10−6)was larger than that of the Fe62Pd30Ni8alloy

(48×10−6). In Fig.5b, the saturation magnetostriction of the Fe62Pd30Ni8alloy SF and ST reveals that adding Ni (8

at.%) into the Fe70_−xPd30Nix FSM alloys lowers the

satu-ration magnetostriction. The reason is unclear. Figure5c, d reveals the saturation magnetostriction of the Fe66Pd30Ni4

and Fe62Pd30Ni8 alloys SF, ST, and then aged at 5 HTA.

The aged Fe66Pd30Ni4alloy had a lower saturation

magne-tostriction (19.3×10−6)and a drastically decreased magne-tostrictive susceptibility (λs/H ), shown in Fig.5c. The Fig. 5 The linear saturation

magnetostriction λs(×10−6)vs.

magnetic field (H )λs− H

curves measured at room temperature: a Fe66Pd30Ni4and

b Fe62Pd30Ni8FSM alloys SF

to a∼38 % reduction in

thickness, ST at 1100◦C for 8 h, and quenched in ice brine;

c Fe66Pd30Ni4and

d Fe62Pd30Ni8alloys SF then

post ST and aged at 5 HTA

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Fig. 6 Magnetization (M) vs.

magnetic field (H ) curves measured at temperature of 300 K (RT) for the Fe66Pd30Ni4

and Fe62Pd30Ni8alloys

strain-forged (SF) to a∼38 % reduction in thickness, ST at 1100◦C for 8 h, and the SF alloys post ST then aged at 5 HTA

reason is that 5 HTA aging treatment caused the decom-position of the L10+ L1m twin phases into stoichiometric

L10+ L1m+ αbctstructures, as shown in Fig.2. However,

the aged Fe62Pd30Ni8alloy maintained a stable saturation

magnetostriction (48.7×10−6), shown in Fig.5d. The rea-son is that the decomposition of L10+L1mtwin phases into

stoichiometric L10+ L1m+ αbctstructures was suppressed

after the 5-HTA aging treatment, as confirmed by the TEM images shown in Fig.4.

3.3 Magnetic Property of the SF Fe66Pd30Ni4

and Fe62Pd30Ni8FSM Alloys ST and the SF Alloys Post

ST Then Aged at 5 HTA

The mass magnetization (M) vs. magnetic field (H ) curves of the SF Fe66Pd30Ni4and Fe62Pd30Ni8alloys ST and the

SF alloys post ST then aged at 5 HTA are shown in Fig.6, where the magnetic applied field is parallel to the sample’s length (i.e., in plane). As shown in Fig.6, the M− H curves of the SF Fe66Pd30Ni4and Fe62Pd30Ni8alloys post ST were

magnetically soft, with a small coercivity Hc < 50 Oe.

The hysteresis loops revealed that the saturation magneti-zation (Ms) of the magnetic moment per unit mass was Ms= 140.15 (emu/g) for the Fe66Pd30Ni4and Ms= 142.90

(emu/g) for the Fe62Pd30Ni8alloys, and both samples had

an anisotropy applied field Ha = 3 kOe, respectively. The

value of the magnetocrystalline anisotropy energy constant (Ku) from equation Ku = (Ms × Ha × ρ) / 2 [10, 11]

(where Ms is saturation magnetization, Ha is anisotropy

applied field, with theoretical density ρFe66Pd30Ni4= 9.350

(g/cm3); ρFe62Pd30Ni8 = 9.390 (g/cm3)) was obtained

as follows: Ku(Fe66Pd30Ni4) = 1.966 × 106 (ergs/cm3);

Ku(Fe62Pd30Ni8)= 2.013 × 106(ergs/cm3). Shown in Fig.6

indicates the M− H curve of the SF Fe66Pd30Ni4alloy after

ST and aging at 5 HTA, revealing it to be magnetically hard, with a coercivity Hc = 500–600 Oe and Ms = 137.13

(emu/g). However, Fig.6reveals that the SF Fe62Pd30Ni8

alloy post ST and aging at 5 HTA was still magnetically soft, with coercivity Hc<50 Oe and Ms = 139.57 (emu/g).

The value of the magnetocrystalline anisotropy energy con-stant (Ku)of the 5-HTA aged Fe66Pd30Ni4and Fe62Pd30Ni8

alloys was obtained as follows: Ku(Fe66Pd30Ni4) = 1.923 ×

106(ergs/cm3); Ku(Fe62Pd30Ni8) = 1.966 × 106(ergs/cm3).

From the hysteresis loop data, it was found that the mag-netic properties of the SF Fe62Pd30Ni8alloy ST and aged

at 5 HTA had a higher saturated magnetization (Ms)and a

higher magnetocrystalline anisotropy energy constant (Ku).

This was due to the low Ni content of the Fe66Pd30Ni4

alloy, which allowed L10+ L1mtwin phases to decompose

into the stoichiometric L10+ L1m+ αbctstructures during

aging at 5 HTA, while in the high Ni content Fe62Pd30Ni8

alloy, they did not. The change in magnetic properties was related to the decomposition of the L10 + L1m

marten-sitic twin phase, as confirmed here by the M − H curves and saturation magnetostriction test in association with microstructure observation, as confirmed by TEM, shown in Figs.2and4.

4 Conclusions

A higher Ni content in the Fe70−xPd30Nix (Nix = 8 at.%)

alloy increases the saturation magnetostriction, satura-tion magnetizasatura-tion (Ms), and magnetocrystalline anisotropy

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energy constant (Ku)after the alloy is SF, ST, and aged at 5

HTA. The reason is that the decomposition of the L10+L1m

twin phases into the stoichiometric L10 + L1m + αbct

structures is suppressed after the 5-HTA aging treatment; therefore, the high Ni content Fe62Pd30Ni8alloy is suitable

for application in high-temperature (T < 500 ◦C) and high-frequency environments.

TEM selected area diffraction pattern (SADP) demon-strated that the orientation relationships of stoichiometric L10+ L1m+ αbct structures in the SF Fe66Pd30Ni4 FSM

alloy after ST and aging at 5 HTA are[ 1 0 1]L10//

[ 4 1 5]L1m//[2 3 3]αbct.

Acknowledgments The author would like to express his sincere appreciation to the Ministry of Science and Technology of the Repub-lic of China for supporting this study (under Grant-in-Aid for MOST 101-2221-E-151-008). The author also wishes to thank Hsueh-Yen Yao of the National Cheng Kung University for his help in the TEM operation.

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