Phase transformation kinetics and microstructural evolution of MnAl permanent magnet alloys

Accepted Manuscript Phase transformation kinetics and microstructural evolution of MnAl permanent magnet alloys Wei Lu, Junchao Niu, Taolei Wang, Kada Xia, Zhen Xiang, Yiming Song, Zhongliang Mi, Weifang Zhang, Wei Tian, Yan Yan PII:

S0925-8388(16)32018-7

DOI:

10.1016/j.jallcom.2016.06.285

Reference:

JALCOM 38154

To appear in:

Journal of Alloys and Compounds

Received Date: 16 May 2016 Revised Date:

27 June 2016

Accepted Date: 29 June 2016

Please cite this article as: W. Lu, J. Niu, T. Wang, K. Xia, Z. Xiang, Y. Song, Z. Mi, W. Zhang, W. Tian, Y. Yan, Phase transformation kinetics and microstructural evolution of MnAl permanent magnet alloys, Journal of Alloys and Compounds (2016), doi: 10.1016/j.jallcom.2016.06.285. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

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Graphical Abstract

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Phase transformation kinetics and microstructural evolution of MnAl permanent magnet alloys Wei Lu1*, Junchao Niu1, Taolei Wang1, Kada Xia1, Zhen Xiang1, Yiming Song1*, Zhongliang Mi2, Weifang Zhang2, Wei Tian2, Yan Yan2 School of Materials Science and Engineering, Tongji University, Shanghai 200092 China

2

Department of Criminal Investigations, Public Security Bureau of Shanghai, Shanghai, 200083

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1

China

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* Corresponding authors. W. Lu: [email protected]; Y. Song: [email protected]

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Abstract:

In order to provide a better understanding about the formation of ferromagnetic τ phase in MnAl Rare-earth-free permanent magnet alloy, the phase transformation kinetics of τ phase has been analyzed for Mn55Al45 alloy. The local activation energy Ec(α) was introduced to explore the change of energy in the process of phases transformation and its value is closely related to the

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transformation volume fraction. The local activation energy is low (about 175 kJ/mol) in the process of nucleation. However, the threshold value of energy will increase slowly and approaches a constant value (about 188 kJ/mol) when the grain-boundary-nucleated τ phase steadily grows into the parent ε-phase grains. With the exception of anomalous stage, the local Avrami exponents

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fall in the range of 2.0-4.0, which implies that two- and three-dimensional nucleation and grain growth controlled by diffusion play a crucial role in the whole transformation process. Meanwhile,

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transmission electron microscopy (TEM) morphology and the results of transformation kinetics are accordance with each other, indicating that the variable local activation energies Ec(α) and local Avrami exponents n(α) are applicable and correct in describing the τ-phase transformation behavior.

1. Introduction

MnAl alloys with compositions slightly on the Mn-rich side of the equiatomic composition exhibit a ferromagnetic τ phase (P4/mmm, also known as L10 phase), which have unique characteristics such as high magnetocrystalline anisotropy and high maximum magnetic energy density [1-4]. As one of the most promising Rare-Earth-free permanent magnets, τ-phase transformation of MnAl

ACCEPTED MANUSCRIPT alloys were investigated to further optimize the microstructure and magnetic properties. Based on the Mn-Al binary phase diagram, the tetragonal τ phase derives from the high-temperature parent ε-phase upon slow cooling from the ε-phase field or under the appropriate thermal treatments after retaining the parent phase at room temperature. The conventional wisdom is that the τ phase is the

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product of a displacive or martensitic transformation following an orthorhombic ε’ phases (B19) ordering within the ε-phase (ε→ε’→τ)[5]. A simple shear of the ε’ phase produces the τ phase with an L10 structure. Early transmission electron microscopy (TEM) studies also provided evidence for such a sequential reaction of chemical ordering by diffusion followed by a displacive shear mode [5].

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The transformation of ε-phase to metastable ε’ phase occurs through a competing chemical ordering at temperatures below 580 . That is to say, the metastable ε’ phase can’t exist in MnAl system at

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high temperature. However, the metastable τ phase typically appears below 800℃, indicating that the τ phase forms without prior ε (hcp) → ε’ (B19) ordering at temperatures of 580℃-800℃[6]. Muller et al summarized the metallurgy of the τ pahse in MnAl-based alloys and pointed out that the

formation of τ phase might not involve the shear mechanism within the ε’ phase [7]. Metallographic studies and high-resolution electron microscopy (HREM) observations of the

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MnAl-base alloy indicated that the τ-phase formation, occurred primarily through a compositionally invariant nucleation and growth mechanism, was akin to the so-called ‘’massive transformation’’. As a distinct genre of solid-state transformation, massive transformation is generally defined as a compositionally invariant nucleation and growth process which involves a

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change in crystal structure [8]. It is considered that the parent phase transforms into the new phase

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through the migration of interphase interfaces and the rate of transformation is controlled by interfacial boundary diffusion processes. The diffusional massive mode has been made a deep research on the microstructural development of TiAl-based alloys

[9-10]

. Later, Hoydick et al.[11-12]

reported the occurrence of a diffusional nucleation and growth process in MnAl system, in keeping with the characteristic of massive transformation. According to the diffusional massive mode, the formation of the ferromagnetic τ phase involves nucleation almost exclusively at prior parent ε-phase grain boundaries followed by growth via motion of mostly incoherent heterophase interface segments [8], which means that the τ-phase originated from the pre-existing parent phase boundaries and essentially is independent of the ε’ (B19) ordering. Previous work about the formation of massive product phases at grain-boundary mainly

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[3,7,8,13]

. The application of

transformation kinetics in the formation of τ-phase has not been deeply studied. In fact, the diffusion-controlled phase transformation behavior can be well investigated through thermomechanical analysis and transformation kinetics

[14]

. Owing to the critical role of atoms

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diffusion near interphase boundary, the study of τ-phase transformation kinetics can also provide a better understanding of τ phase formation in MnAl alloys. In this study, the influence of annealing temperatures and heating rates on the phase transformation were systematically explored. And on this basis we introduce the concept of local activation energies Ec(α) and transformation

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parameters such as Avrami exponent n to analyze early transition process of the ferromagnetic τ-phase. Kinetic parameters of τ-phase transformation can be obtained from differential scanning

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calorimetry (DSC) results at various heating rates. The transmission electron microscopy (TEM) morphology and the calculations results of transformation kinetics are accordance with each other, indicating that the variable local activation energies Ec(α) and local Avrami exponents n(α) are applicable and correct in describing the τ-phase transformation behavior. In addition, the study of phase transformation kinetics and microstructural evolution can contribute to the unraveling of

2. Experimental

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many controversial issues about the nucleation and growth behavior in the formation of τ-phase.

Mn55Al45 alloy ingots were prepared by arc melting under an argon atmosphere using

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high-purity elemental materials (99.98 pct Mn, 99.99 pct Al). The ribbons were obtained by ejecting the melt from a quartz tube onto the surface of a rotating copper wheel under argon

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atmosphere. The rotating rate 30m/s was selected to form the pure parent ε-phase during the fabricating process. In order to investigate the ferromagnetic τ-phase transformation process , the as-spun ribbons were annealed at temperatures from 410 to 490

for 10 min in a vacuum furnace

and the heating rate 5 K/min was used in the process of annealing. The crystallographic structure was characterized by X-ray diffraction (XRD) with Cu Kα radiation. The ferromagnetic phase transformation kinetics of as-spun alloy ribbons was investigated with a NETZSCHSTA449C TG/DSC instrument using a continuous heating regime with heating rates from 5 to 40℃/min. A N2 gas flow was used for the ambient atmosphere. The electron transparent foils for TEM experiments have been prepared by electropolishing with a twin-jet polisher at room temperature. 3. Results and discussions

ACCEPTED MANUSCRIPT In order to induce parent phase transformation and get high purity τ-phase, appropriate annealing condition was determined with differential scanning calorimetry (DSC). Fig.1 shows the DSC curves of the as-prepared alloy ribbon taken at 5-40 K/min. The ferromagnetic phase transformation process of Mn55Al45 alloy ribbon at different heating rates exhibits a single

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exothermal peak in the DSC curves which are normalized with sample weight. The first exothermal peak of the 5 K/min curve is in the range of 700K-710K and the peak temperature gradually moves to right with increasing heating rates until 40 K/min. This means that a higher heating rate will delay the optimum transformation temperature in the formation of τ phase.

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Besides, several annealing temperatures for the melt-spun ribbons are chosen according to these peaks in DSC curve. The XRD patterns of the ribbons annealed at these temperatures are shown in

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Figure 2. It is found that the as-spun ribbon consisted of mostly ε-phase and a small amount of τ-phase. After annealing treatment, the parent ε phase was effectively transformed to τ phase. Depending on the annealing temperature and time, the annealed ribbons were mainly composed of three phases: τ phase, γ2 phase, and β phase. When the ribbons were annealed at 430

for 10min,

the ribbons still contained ε phase. This implies that the ε phase did not transformed to τ phase

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completely. Increasing the annealing temperature will promote the ε→τ transformation. It is noted that only pure τ phase was obtained in the ribbon annealed at 450

for 10 min, implying that the

sample annealed at the optimum temperature might contain more transition state of ε→τ

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transformation. In order to observe the electron microscopy morphology involving τ phase nucleation and growth, those alloy ribbons should be studied systematically. Meanwhile, the

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optimum annealing condition is consistent with the peak temperature of 5 K/min DSC curve, which caused by the fact that the heating rate 5 K/min was used in the process of annealing. The continuous increase of annealing temperature will lead to the decomposition of the metastable τ-phase and the equilibrium impure phases appeared in the ribbons. Just as shown in Fig.2, annealing at temperature higher than 470

and 490

resulted in a large fraction of non-magnetic

γ2 and β phases, and the volume fraction of τ phase in MnAl alloy ribbon decreases obviously. From the DSC curves measured at different heating rates in Fig.1, different kinetic parameters of the τ-phase transformation in MnAl alloys can be obtained according to the classical transformation dynamics theory[14,15], such as activation energy of phase transformation .Generally, activation energy is defined as the threshold value of energy above which the fluctuation of energy

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variable activation energy and reflect the change of nucleation and growth behavior in the formation of τ-phase, a local activation energy Ec(α) is introduced, which presents the activation energy at a stage when the transformed volume fraction is α. Considering the difference of alloy systems, an isoconversional method without assuming the kinetic model function is used to

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calculate the value of Ec(α)[16]. This model-free method is expressed in the form

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 d ln φ  Ea  d (1 T ) = − R  α

where R is gas constant, T andφare the temperatures and specific heat flow (w/g) corresponding to fixed values of α from experiments at different heating rates β, respectively. The curve of lnφ vs. 1/T at a certain value of α can be plotted and Ec(α) is obtained from the slope of the straight line. Fig.3 shows the relationship between transformation volume fraction and local activation

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energy of MnAl alloys. In the beginning, the local activation energies for τ-phase transformation are 177±2KJ/mol(α=0.1) and 184±4 KJ/mol(α=0.2). It is obvious that the local activation energy increases slowly and approaches a constant value(about 188 ± 2KJ/mol) when the

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transformation volume fraction falls in the range of 0.3-0.9. However, the local activation energy increases abnormally at the last stage of τ-phase transformation (α>0.9). At the initial stage of

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phase transformation in MnAl alloys, the activation energy is expected to be dominated by the nucleation process for a diffusional massive transformation. In fact, the chemical surrounding of forming massive τ-product contains high-density defects, such as twin boundaries, micro-twins, stacking faults and dislocations[17,18]. The diffusional nucleation preferentially forms at the grain boundaries and defect positions, reducing the local activation energy at the beginning of transformation. With the increase of transformation volume fraction (α=0.3-0.9), the grain-boundary nucleated τ phase steadily grows into the parent ε-phase grains, which is the main reason for the constant local activation energy during the period of τ-phase growth. When the transformation volume fraction reaches 0.9 and higher, there may exist the decomposition of the metastable τ-phase according to the reaction equation τ γ2+β, leading to the local activation

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2KJ/mol, which is lower than Ozawa model 191±2KJ/mol. Considering the value of local activation energy, it can be seen that Ec(α) falls in between Kissinger model and Ozawa model, but there are no significant difference between the three.

The diffusion-controlled grain growth and transformation can be most frequently described

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by Avrami exponent. This kinetic parameter is related to the growth mechanism of new phases and has proved to be significant in describing the transformation behavior[20,21]. Considering the

which is expressed as follows [22]:

n(α ) =

− R∂ ln  − ln (1 − α )  E (α )∂ (1/ T )

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non-isothermal transformation and local activation energy, we introduce a local Avrami exponent,

The double logarithmic plots ln(-ln(1-α)) vs. 1/T were obtained from the DSC data for the τ-phase

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transformation. Owning to the variable local activation energy Ec(α), the local Avrami exponents n(α) at heating rate of 5K/min were calculated. Fig.4 shows the value of local Avrami exponent of MnAl alloys during the τ-phase transformation. It is obvious that the Avrami exponent varies

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significantly with α in the whole process of transformation. At the beginning, n(α) is about 8.5. With the transformation of parent phase, the value of n(α) decreases sharply from 8.5 to about 4 when α resches 0.03. In this stage, the calculated local Avrami exponent n(α) is larger than 4.0,

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which is anomalous according to transformation theory [21] and has no physical meaning. This may be caused by the uncertainties and complexity of initial stage(α 0.03

and the DSC date rely

on the choice of baseline. Therefore, the anomalous value for α

0.03 is not taken into

consideration. When the transformation volume fraction α is in the range of 0.03-1.0, the Avrami exponent decreases slowly to about 2.0. It means that two- and three-dimensional nucleation and grain growth controlled by diffusion at decreasing nucleation rate are dominant for τ-phase transformation in this stage

[22]

. The n(α)-α dependence curves obtained at other heating rates are

similar to this one. The τ-phase nucleation and growth mechanism obtained from the local Avrami exponents

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for 10 min.

The three-dimensional nucleation characteristic that shown in Fig.5(a) is consistent with the thermodynamics calculation. It is obvious that the heterogeneous nucleation of the τ-phase occurs

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only at the grain boundaries of the parent ε-phase after isothermal treatment. Fig.5(b) clearly reveals the early τ-phase nuclei at the grain boundaries and the faceting of the interphase interfaces. These observations are also consistent with the idea that the τ-phase nuclei approximate a single-faceted nucleus [23]. Meanwhile, it can be observed that the faceted τ phase grows only into

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one side of the ε-phase grain during the early transformation stages, implying that the growth mode is accomplished with the migration of incoherent heterophase interfaces by random atomic

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diffusion and attachment across the growth interface [8,24]. 4. Conclusions

In summary, the τ-phase transformation kinetics of MnAl alloys has been analyzed by non-isothermal DSC measurements, providing a better understanding of τ-phase phase formation. The local activation energy Ec(α) was introduced to explore the process of phases transformation.

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At initial stage of τ-phase transformation(α 0.2), the diffusional nucleation preferentially forms at grain boundaries and defect positions, resulting in that the value of Ec(α) is lower than that of the growth stage . The abnormal increase at the final stage may be attributed to the decomposition

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of the metastable τ-phase, which is consistent with the evolution of phase composition observed in XRD patterns. The variation of local Avrami exponent with transformation fraction demonstrates

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the mechanism of τ-phase nucleation and growth. With the exception of anomalous stage, the local Avrami exponents fall in the range of 2.0-4.0, indicating that two- and three-dimensional nucleation and grain growth controlled by diffusion play a crucial role in the whole transformation process. Meanwhile, The experimental results showed that the TEM morphology and the calculations of transformation kinetics are accordance with each other, indicating that the variable local activation energies Ec(α) and local Avrami exponents n(α) are applicable and correct in describing the τ-phase transformation behavior.

Acknowledgement This work was supported by the National Natural Science Foundation of China (Grant No 50901052), the ‘Morning Star’ project (Grant No 14QA1403600) supported by the Science and

ACCEPTED MANUSCRIPT Technological Commission of Shanghai, the Fundamental Research Funds for the Central Universities and Key Laboratory of Impression Evidence Examination and Identification Technology, Ministry of Public Security, PRC (Grant No HJJY2014010).

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Figure caption: Fig.1 Non-isothermal DSC curves for the phase transformation of MnAl alloys Fig.2 XRD patterns for as-spun ribbons before and after heat treatment

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Fig.3 The dependence of the local activation energy Ec(α) on the transformation fraction Fig.4 The average local Avrami exponent n(α) as a function of transformation fraction a at heating rate of 5K/min.

Fig.5 (a) TEM bright-field micrograph depicting the characteristics of three-dimensional nucleation

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controlled by diffusion at the grain boundaries of the parent ε-phase

(b) Bright-field transmission electron microscopy showing that the growing τ-phase with faceted

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interphase interfaces grows only into one of the ε-phase grains during the early transformation stages

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Highlights 1)

The phase transformation kinetics of τ phase has been analyzed in Mn55Al45 alloy;

2)

Two- and three-dimensional nucleation and grain growth controlled by diffusion play a crucial role in the ε to τ phase transformation process;

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Local activation energies and local Avrami exponents are applicable in describing the τ phase

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transformation behavior.

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3)