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Article

Effect of Mo Alloying on the Precipitation Behavior of B2 Nano-Particles in Fe-Mn-Al-Ni Shape Memory Alloys

by
Liqiu Yong
1,
Yang Zuo
1,
Lixin Sun
1,
Huabei Peng
1,*,
Xuguang An
2,
Hui Wang
2 and
Yuhua Wen
1
1
School of Mechanical Engineering, Sichuan University, Chengdu 610065, China
2
Interdisciplinary Materials Research Center, Institute for Advanced Study, Chengdu University, Chengdu 610106, China
*
Author to whom correspondence should be addressed.
Metals 2022, 12(2), 261; https://doi.org/10.3390/met12020261
Submission received: 29 December 2021 / Revised: 24 January 2022 / Accepted: 29 January 2022 / Published: 30 January 2022
(This article belongs to the Special Issue Novel Shape Memory Alloys)
Browse Figures
Versions Notes

Abstract

:
In Fe-Mn-Al-Ni shape memory alloys, the stabilization of superelasticity would be affected by the undesired precipitation of B2 nano-particles during natural aging. In order to solve this problem, the effect of Mo alloying on the precipitation behavior of B2 nano-particles during the cooling and natural aging processes was performed by scanning electron microscope, transmission electron microscope and Vickers microhardness test in two Fe-Mn-Al-Ni-Mo shape memory alloys. The results showed that the formation of γ phase was completely suppressed after 15 °C and 80 °C water quenching as well as air cooling. However, B2 nano-particles were still precipitated after the three cooling processes, and their sizes and misfits increased with decreasing the cooling rates. In addition, the Vickers hardness increased after natural aging for 338 days, which indicated that it is not viable to inhibit the precipitation of B2 nano-particles during natural aging by Mo alloying in the Fe-Mn-Al-Ni shape memory alloys.

1. Introduction

Shape memory alloys (SMAs) have attracted much attention as a class of novel smart alloys owing to their shape memory effect and superelasticity [1]. Among a large number of SMAs [2,3,4,5,6,7,8,9,10,11,12,13,14,15,16], Fe-based SMAs are expected to replace Ni-Ti SMAs for large-scale industrial applications due to their low cost and high workability [17,18,19]. In particular, Fe-Mn-Al-Ni-based SMAs have temperature-insensitive superelastic stress and exhibit superelasticity over a wide temperature range [20,21,22]. Owing to this unique feature, they are suitable for working in extreme environments such as the moon and Mars. For the Fe-Mn-Al-Ni-based SMAs, B2 nano-particles play a crucial role in obtaining the superelasticity. A number of studies reported that the best superelasticity was obtained when the diameter of B2 nano-particles was about 10 nm, and the superelasticity decreased when the diameters deviated from this value [22,23,24]. In this case, aging treatment at 200 °C is commonly performed to obtain the B2 nano-particles with the optimum diameter of 10 nm in the Fe-Mn-Al-Ni-based SMAs. However, the optimum size of B2 nano-particles is unstable and could be further changed by the effect of natural aging in the aged Fe-Mn-Al-Ni-based SMAs. Ozcan et al. [24] found that the average size of B2 precipitates increased from 5 to 7 nm and created 5% reversible superelasticity after room temperature aging for 30 days in a Fe43.5Mn34Al15Ni7.5 (at.%) SMA. This further evolution of B2 nano-particles during natural aging would affect the stabilization of superelasticity in Fe-Mn-Al-Ni SMAs. However, there are no reports on effective ways to address this issue.
Based on the Lifshitz–Slyozov–Wagner (LSW) theory, the diffusivity of partitioning elements is one of the key factors to control the growth of nano-particles. Accordingly, it may be able to suppress the precipitation of nano-particles by impeding the diffusion of constituent elements. Jiang et al. [25] reported that high-density B2 nano-particles with extremely fine size (2.7 ± 0.2 nm) were precipitated owing to the addition of Mo in a novel Fe-18Ni-3Al-4Mo-0.8Nb-0.08C-0.01B (wt.%) maraging steel. They explained that Mo impedes the element diffusion of B2 nano-particles, as its diffusion rate is lower than that of the other constituent elements. Meanwhile, the elastic misfit energy (driving force for coarsening) is reduced by lowering the lattice misfit between matrix and B2 nano-particles owing to the partitioning of Mo in the α matrix. In the Fe-Mn-Al-Ni SMAs, Mo alloying may also be viable to suppress the precipitation of B2 nano-particles during natural aging. Inspired by this, we designed two Fe-Mn-Al-Ni-Mo SMAs with a Mo content of about 2.3 at.% to investigate the effect of Mo alloying on the precipitation behavior of B2 nano-particles. This Mo content is consistent with that of the Fe-Ni-Al-Mo-Nb-C-B maraging steel, which has been proved to suppress the precipitation and growth of B2 nano-particles significantly. The results showed that the γ phase was completely inhibited even after air cooling, but the precipitation behavior of B2 nano-particles was not suppressed by Mo alloying during both cooling and natural aging processes in the Fe-Mn-Al-Ni-Mo SMAs.

2. Materials and Methods

The Fe-Mn-Al-Ni-Mo SMAs ingots were prepared by vacuum-induction melting under an argon atmosphere, using high-purity iron, manganese, silicon, aluminum, nickel and molybdenum. Then, the ingots were homogenized at 1100 °C for 15 h and then hot forged into bars with 20 mm diameter. The specimens of 10 × 5 × 2 mm3 were cut from the bars by an electro-discharged machine. In order to investigate the effect of cooling rates on the precipitation behavior of B2 nano-particles, the specimens were solution treated at 1150 °C for 30 min followed by water quenching at 15 °C (15 °C WQ), water quenching at 80 °C (80 °C WQ) and air cooling at room temperature (AC), respectively. Among them, the 15 °C water is uniformly used for the general water quenching at room temperature because the precipitation of B2 nano-particles is rapid and sensitive to the change in cooling time. The cooling rate of 80 °C WQ is lower than that of 15 °C WQ. Meanwhile, the 80 °C WQ was also utilized to precisely precipitate a small amount of ductile γ phase along the grain boundary, which could suppress the crack initiation in the polycrystalline Fe-Mn-Al-Ni SMAs [26]. Chemical compositions of two Fe-Mn-Al-Ni-Mo SMAs are given in Table 1. According to the difference in Ni contents, the two alloys were referred to as 6Ni2Mo and 8Ni2Mo alloys, respectively.
For microstructure observation, specimens were mechanically ground and then electrolytically polished in an alcoholic solution of 10% perchloric acid for 30 s with a constant voltage of 20 V. Microstructure observation was then carried out by a Phenom Pro scanning electron microscope (SEM) (Thermo Fisher Scientific, Waltham, MA, USA). Furthermore, the precipitation behavior of B2 nano-particles was characterized by an FEI Talos F200X transmission electron microscope (TEM) (Thermo Fisher Scientific, Waltham, MA, USA). The average diameter of B2 nano-particles was obtained by counting 150 nano-particles at least. It was reported that the precipitation of nano-particles is correlated with the hardness values of materials [23,27,28,29]. Therefore, a Vickers hardness test with a load of 9.8 N was conducted to characterize the precipitation behavior of B2 nano-particles during natural aging. The Vickers hardness of each specimen was measured at least 5 times for data reliability. Besides, crystal orientation also affects the hardness of the material [30]. In order to avoid the effect of different crystal orientations on the Vickers hardness, all specimens for natural aging were cut from one large single crystal, which was fabricated by cyclic heat treatment [31,32].

3. Results

3.1. Microstructures of Fe-Mn-Al-Ni-Mo SMAs

Figure 1 gives the SEM micrographs of solution-treated 6Ni2Mo and 8Ni2Mo alloys subjected to 15 °C WQ or AC. α single phase was observed for both two alloys subjected to 15 °C WQ or AC. It is reported that the γ phase usually forms along the grain boundary of the α phase due to the low cooling rate in the Fe-Mn-Al-Ni SMAs [26]. Our previous studies showed that widmanstatten γ phase was generated in a Fe44.61Mn34.74Al13.38Ni8.27 (at.%) alloy after WQ [33,34]. Furthermore, the γ phase with a fraction of about 40% was precipitated when a Fe44.5Mn34Al14Ni7.5 (at.%) alloy was air-cooled for only 5 s before 15 °C WQ [26]. In contrast to the Fe-Mn-Al-Ni quaternary SMAs, the formation of the γ phase was completely suppressed even under AC in both 6Ni2Mo and 8Ni2Mo alloys. This indicated that Mo alloying would strongly suppress the formation of the γ phase in Fe-Mn-Al-Ni SMAs. A similar low quenching sensitivity of the γ phase was also reported in the Fe-Mn-Al-Ni-Ti SMA [35]. In addition, intragranular cracks were formed after 15 °C WQ but not after AC. On the one hand, the cracking was prone to occur because the precipitation of the ductile γ phase was completely suppressed [26]. On the other hand, tensile thermal stress was generated by the rapid cooling rate of WQ and not effectively reduced owing to the constraints of grain boundaries [36]. As a result, the intragranular cracks were induced by this tensile thermal stress in the 6Ni2Mo and 8Ni2Mo SMAs after 15 °C WQ.

3.2. Precipitation Behavior of B2 Nano-Particles in Fe-Mn-Al-Ni-Mo SMAs

Figure 2 and Figure 3 show the precipitation behavior of B2 nano-particles for solution-treated 6Ni2Mo and 8Ni2Mo alloys followed by different cooling methods. The Electron diffraction patterns and dark field images indicated that B2 nano-particles were precipitated in the α matrix for both 6Ni2Mo and 8Ni2Mo alloys. To be specific, the average diameter of B2 nano-particles was 7.8 nm, 8.7 nm and 25.6 nm for 6Ni2Mo alloys, while it was 11.7 nm, 13.7 nm and 38.9 nm for 8Ni2Mo alloys subjected to 15 °C WQ, 80 °C WQ and AC, respectively. Figure 4 shows the average diameters of B2 nano-particles for the solution-treated 6Ni2Mo and 8Ni2Mo alloys followed by 15 °C WQ, 80 °C WQ or AC. The error bars on the columns represent standard deviations from the mean for the statistical diameters of B2 nano-particles. The average diameter of B2 nano-particles was about three times larger in air-cooled alloys than in the same water-quenched alloys. In addition, the average diameter of B2 nano-particles for 8Ni2Mo alloy was larger than that for 6Ni2Mo alloy under the same cooling method. Because Ni is one of the forming elements of B2 nano-particles, higher Ni content would promote their growth.
Figure 5 and Figure 6 show HRTEM images, Fast Fourier transform (FFT) images and FFT filtered images for the solution-treated 6Ni2Mo and 8Ni2Mo alloys subjected to different cooling methods. The morphology of B2 nano-particles deviated from sphericity to rounded rectangle as the cooling rates slowed down in both two alloys. The lattice constants of α matrix and B2 nano-particles were quantitatively determined by the Fast Fourier transform of HRTEM images, as shown in Table 2. The lattice constants of 6Ni2Mo alloy were larger than that of 8Ni2Mo alloy under the same cooling methods. Meanwhile, the lattice constants of the α matrix were consistently bigger than that of the B2 phase, and their misfits gradually increased with reducing cooling rates for both two alloys. Note that the misfits were calculated as follows [25]:
2 | a α − a B 2 | a α + a B 2 × 100 %
where aα and aB2 is the lattice constant of the α matrix and B2 nano-particles, respectively.

3.3. Effect of Aging Treatment on Vickers Hardness in Fe-Mn-Al-Ni-Mo SMAs

Figure 7 shows the effect of natural aging for 338 days on the microhardness of 15 °C water-quenched 6Ni2Mo and 8Ni2Mo alloys subjected to aging at 200 °C at different times. The Vickers hardness was 414 and 423 in the solution-treated 6Ni2Mo and 8Ni2Mo alloy without aging process, respectively, while it increased to 428 and 429 after natural aging for 338 days. Besides, the Vickers hardness increased and then decreased as the aging time at 200 °C was prolonged, while the peak hardness was obtained at 7 h in the 6Ni2Mo alloy. For the 8Ni2Mo alloy, however, the Vickers hardness drastically increased for 1 h and then decreased until 7 h, followed by a further increase. After aging at 200 °C at different times, these specimens were then naturally aged for 338 days. Interestingly, all the Vickers hardness of these specimens increased than that before the natural aging.

4. Discussion

It is known that lattice constants are significantly affected by the partitioning of constituent elements [37,38]. The value of atomic radius (R) is RMo >RMn > RAl > RFe > RNi in Fe-Mn-Al-Ni-Mo alloys [39,40]. Since Mo has the largest atomic radius, the crystal lattices of phases would be expanded when their solute atoms are replaced by it. For the Fe-Mn-Al-Ni-based SMAs, it was reported that the α matrix enriches with Fe and Mn while B2 nano-particles enrich with Ni and Al elements [20]. In addition, the Mo atom is found to be expelled from the B2 nano-particles into the α matrix [25,41]. Accordingly, the lattice constant of the α matrix was always larger than that of B2 nano-particles due to its enrichment of Mn and Mo with a larger atomic radius (Table 2). As the cooling rates decreased, more Mo diffused into the α matrix, resulting in a gradual increase in the misfits between the two phases. Nevertheless, the B2 nano-particles were still maintained completely coherent with the α matrix in the diameter range of 7.8~38.9 nm (degrees of misfit <5%).
As mentioned above, B2 nano-particles with small size (about 2.7 ± 0.2 nm) were precipitated by adding Mo in the Fe-Ni-Al-Mo-Nb-C-B alloy [25]. However, the average diameter of B2 nano-particles was still 7.8 nm and 11.7 nm in the solution-treated 6Ni2Mo and 8Ni2Mo alloys subjected to 15 °C WQ, respectively (Figure 2 and Figure 3). Furthermore, the B2 nano-particles coarsened as the cooling rates decreased. These results indicated that the precipitation and growth behavior of B2 nano-particles were not suppressed during the cooling process in the Fe-Mn-Al-Ni-Mo SMAs, which is different from the effect of Mo alloying in the Fe-Ni-Al-Mo-Nb-C-B alloy. It could be explained by the large misfits between the α matrix and B2 nano-particles in the Fe-Mn-Al-Ni-Mo alloys. Compared to the low lattice misfit (0.03 ± 0.04%) of B2 nano-particles in the Fe-Ni-Al-Mo-Nb-C-B alloy, the larger misfits of B2 nano-particles in the Fe-Mn-Al-Ni-Mo alloys lead to higher elastic mismatch energy. This elastic mismatch energy could be the driving force for precipitates growth and eventually coarsens the B2 nano-particles.
The Vickers hardness evolutions are significantly different between the 6Ni2Mo and 8Ni2Mo alloys (Figure 7). Since the identical heat treatments were employed, it should be attributed to the difference in compositions of the two alloys. The precipitation behaviors vary with compositions and directly affect the Vickers hardness. First, the average diameter of B2 nano-particles is larger in the 8Ni2Mo alloy than in the 6Ni2Mo alloy after 15 °C WQ owing to the higher Ni content (Figure 2 and Figure 3). Accordingly, the initial Vickers hardness is higher in the 8Ni2Mo alloy (423 HV) than in the 6Ni2Mo alloy (414 HV) after 15 °C WQ. When beginning to age at 200 °C, the B2 nano-particles decompose from the supersaturated α matrix and raise their volume fractions. The higher the volume fraction of precipitates in the matrix, the higher the hardness value [29]. Meanwhile, the coarsening of B2 nano-particles increases the difficulty of dislocation movement, which also increases the hardness at the beginning of 200 °C aging. Compared with the 6Ni2Mo alloy, the precipitation of B2 nano-particles is promoted by the higher Ni content in the 8Ni2Mo alloy, and thus its Vickers hardness increases more rapidly at this stage. As the aging time is prolonged, the volume fraction of B2 nano-particles would not continue to increase, but the B2 nano-particles keep coarsening to reduce the interface energy. In this case, the density of B2 nano-particles decreases, and the interparticle spacing becomes larger. The dislocation movement becomes easier, and the Vickers hardness decreases according to the classic Orowan theory [28]. It is worth noting that a further increase in hardness appeared after aging more than 7 h at 200 °C in the 8Ni2Mo alloy. This increase in hardness might be related to the precipitation of new phases such as Ni3Mo or β-Mn. The Ni3Mo nano-phase is a common strengthening phase in the iron-nickel-based maraging steels [42,43]. As Ni is one of the forming elements of Ni3Mo, its higher content may facilitate the precipitation of Ni3Mo nano-phase in the 8Ni2Mo alloy. However, it is reported that the precipitation of Ni3Mo nano-phase usually occurs at temperatures above 500 °C and high Ni contents (e.g., 18 wt.%) in the iron-nickel-based maraging steels, while our aging temperature and Ni contents are both lower than that. Besides, our earlier study [33] discovered that β-Mn nano-particles were precipitated in a Fe43.61Mn34.74Al13.38Ni8.27 alloy (at.%) after air cooling, which also might be precipitated in the aged 8Ni2Mo alloy. Obviously, the precipitation behavior of nano-phase varies with the composition, and the type of new phase cannot be confirmed only by the Vickers hardness. Thus, more investigations would be conducted to clarify the precipitation behavior of B2 nano-particles and identify the new phase in the future.
After natural aging for 338 days, the Vickers hardness increased in both 15 °C water-quenched 6Ni2Mo and 8Ni2Mo alloys (Figure 7). This indicated that the B2 nano-particles were still precipitated and grew in the 15 °C water-quenched alloys during the natural aging. Ozcan et al. also reported that the average diameter of B2 nano-particles increased from 5 to 7 nm after natural aging for 30 days in a water-quenched Fe43.5Mn34Al15Ni7.5 (at.%) SMA [24]. Interestingly, the increase in hardness caused by natural aging also occurred at the stage where the hardness decreased with increasing aging time at 200 °C. It means that this increase in hardness is not only induced by the coarsening of B2 nano-particles but also may be related to the further partitioning of elements in the α matrix during natural aging [38]. As discussed above, the partitioning of Mo leads to an increase in the misfits of B2 nano-particles, which intensifies the elastic strain field around them. In this case, the dislocation motion would be impeded by the strengthened strain field. Eventually, all the Vickers hardness values are elevated after the natural aging. This increase in Vickers hardness indicated that Mo alloying could not suppress the precipitation of B2 nano-particles during natural aging. Therefore, it is not viable to improve the stabilization of superelasticity by Mo alloying during the natural aging in the Fe-Mn-Al-Ni SMAs.

5. Conclusions

In this study, the effect of Mo alloying on the precipitation behavior of B2 nano-particles during cooling and natural aging was investigated by the TEM and Vickers hardness test in two Fe-Mn-Al-Ni-Mo SMAs. The main conclusions are as follows:
(1)
The formation of the γ phase was completely suppressed even under AC, but the intragranular cracks were generated after 15 °C WQ in both two Fe-Mn-Al-Ni-Mo SMAs;
(2)
B2 nano-particles were precipitated during the cooling process, and their sizes increased as the cooling rates decreased in both two Fe-Mn-Al-Ni-Mo SMAs;
(3)
The misfits between B2 nano-particles and α matrix increased with coarsening the B2 nano-particles in both two Fe-Mn-Al-Ni-Mo SMAs;
(4)
After natural aging for 338 days, all the Vickers microhardness increased in the 15 °C water-quenched 6Ni2Mo and 8Ni2Mo alloys subjected to aging at 200 °C for different times. It indicated that the precipitation of B2 nano-particles during natural aging was not suppressed by the Mo alloying in the Fe-Mn-Al-Ni SMAs.

Author Contributions

Conceptualization, L.Y. and H.P.; Investigation, L.Y., Y.Z., L.S., X.A. and H.W.; Methodology, H.P.; Resources, H.P. and Y.W.; Supervision, H.P.; Validation, L.Y.; Writing—original draft, L.Y.; Writing—review and editing, H.P. All authors have read and agreed to the published version of the manuscript.

Funding

This research was funded by the National Natural Science Foundation of China [No. 51971152] and Sichuan Science and Technology Program [No. 2020YJ0258].

Data Availability Statement

Not applicable.

Conflicts of Interest

The authors declare no conflict of interest.

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Figure 1. SEM micrographs of solution-treated 6Ni2Mo (a) and 8Ni2Mo (b) alloys subjected to 15 °C water quenching (WQ) or air cooling (AC). The magnified images show clean grain boundaries without γ phase.
Figure 1. SEM micrographs of solution-treated 6Ni2Mo (a) and 8Ni2Mo (b) alloys subjected to 15 °C water quenching (WQ) or air cooling (AC). The magnified images show clean grain boundaries without γ phase.
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Figure 2. Electron diffraction patterns, dark field images corresponding to the {100}B2 spot, as well as the size distribution of B2 nano-particles for the solution-treated 6Ni2Mo alloy followed by 15 °C WQ (a), 80 °C WQ (b) or AC (c).
Figure 2. Electron diffraction patterns, dark field images corresponding to the {100}B2 spot, as well as the size distribution of B2 nano-particles for the solution-treated 6Ni2Mo alloy followed by 15 °C WQ (a), 80 °C WQ (b) or AC (c).
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Figure 3. Electron diffraction patterns, dark field images corresponding to the {100}B2 spot, as well as the size distribution of B2 nano-particles for the solution-treated 8Ni2Mo alloy followed by 15 °C WQ (a), 80 °C WQ (b) or AC (c).
Figure 3. Electron diffraction patterns, dark field images corresponding to the {100}B2 spot, as well as the size distribution of B2 nano-particles for the solution-treated 8Ni2Mo alloy followed by 15 °C WQ (a), 80 °C WQ (b) or AC (c).
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Figure 4. Average diameters of B2 nano-particles for the solution-treated 6Ni2Mo and 8Ni2Mo alloys followed by 15 °C WQ, 80 °C WQ or AC.
Figure 4. Average diameters of B2 nano-particles for the solution-treated 6Ni2Mo and 8Ni2Mo alloys followed by 15 °C WQ, 80 °C WQ or AC.
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Figure 5. HRTEM images, FFT images and FFT filtered images for the solution-treated 6Ni2Mo alloy subjected to 15 °C WQ (a), 80 °C WQ (b) or AC (c). Note that the results in the red and green solid box were from the corresponding red and green dashed areas in the HRTEM images, respectively.
Figure 5. HRTEM images, FFT images and FFT filtered images for the solution-treated 6Ni2Mo alloy subjected to 15 °C WQ (a), 80 °C WQ (b) or AC (c). Note that the results in the red and green solid box were from the corresponding red and green dashed areas in the HRTEM images, respectively.
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Figure 6. HRTEM images, FFT images and FFT filtered images for the solution-treated 8Ni2Mo alloy subjected to 15 °C WQ (a), 80 °C WQ (b) or AC (c). Note that the results in the red and green solid box were from the corresponding red and green dashed areas in the HRTEM images, respectively.
Figure 6. HRTEM images, FFT images and FFT filtered images for the solution-treated 8Ni2Mo alloy subjected to 15 °C WQ (a), 80 °C WQ (b) or AC (c). Note that the results in the red and green solid box were from the corresponding red and green dashed areas in the HRTEM images, respectively.
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Figure 7. Vickers hardness of the 15 °C water-quenched 6Ni2Mo (a) and 8Ni2Mo (b) alloys subjected to aging at 200 °C for different times, followed by 338 days natural aging.
Figure 7. Vickers hardness of the 15 °C water-quenched 6Ni2Mo (a) and 8Ni2Mo (b) alloys subjected to aging at 200 °C for different times, followed by 338 days natural aging.
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Table
Table 1. Chemical compositions of Fe-Mn-Al-Ni-Mo SMAs.
Table 1. Chemical compositions of Fe-Mn-Al-Ni-Mo SMAs.
AlloysChemical Compositions (at.%)
FeMnAlNiMo
6Ni2Mo44.336.510.76.22.3
8Ni2Mo43.635.210.97.92.4
Table
Table 2. Average diameters, lattice constants and misfits of B2 nano-particles in Fe-Mn-Al-Ni-Mo SMAs subjected to different cooling methods.
Table 2. Average diameters, lattice constants and misfits of B2 nano-particles in Fe-Mn-Al-Ni-Mo SMAs subjected to different cooling methods.
AlloysCooling MethodsAverage Diameters of B2 Nano-Particles (nm)Lattice Constants (Å)Misfits (%)
α MatrixB2 Nano-Particles
6Ni2Mo15 °C WQ7.82.9422.9320.34
80 °C WQ8.72.9552.9091.57
AC25.62.9782.8703.69
8Ni2Mo15 °C WQ11.72.8852.8810.14
80 °C WQ13.72.9392.8951.51
AC38.92.9062.8143.22
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Yong, L.; Zuo, Y.; Sun, L.; Peng, H.; An, X.; Wang, H.; Wen, Y. Effect of Mo Alloying on the Precipitation Behavior of B2 Nano-Particles in Fe-Mn-Al-Ni Shape Memory Alloys. Metals 2022, 12, 261. https://doi.org/10.3390/met12020261

AMA Style

Yong L, Zuo Y, Sun L, Peng H, An X, Wang H, Wen Y. Effect of Mo Alloying on the Precipitation Behavior of B2 Nano-Particles in Fe-Mn-Al-Ni Shape Memory Alloys. Metals. 2022; 12(2):261. https://doi.org/10.3390/met12020261

Chicago/Turabian Style

Yong, Liqiu, Yang Zuo, Lixin Sun, Huabei Peng, Xuguang An, Hui Wang, and Yuhua Wen. 2022. "Effect of Mo Alloying on the Precipitation Behavior of B2 Nano-Particles in Fe-Mn-Al-Ni Shape Memory Alloys" Metals 12, no. 2: 261. https://doi.org/10.3390/met12020261

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