Microstructure of Materials
Effect of Ball Milling and Spark Plasma Sintering on Microstructure and Properties of Mn–Cu Based Damping Alloy
2021 Volume 62 Issue 9 Pages 1279-1284
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2021 Volume 62 Issue 9 Pages 1279-1284
In the present study, a Mn–Cu-based damping alloy with a nominal composition of Mn–25Cu–2Al was produced by ball milling and spark plasma sintering. The phase composition, microstructure, compactness, micro-hardness, damping capacity of these materials were systematically investigated. Experimental results show that the ball milling mixed particles are mainly composed of α-Mn and γ-Cu. The microstructure of the alloy sintered at 730°C is mainly composed of γ-MnCu, γ-CuMn and α-Mn. With the increase of sintering temperature, γ-CuMn and α-Mn gradually disappear by the mutual diffusion of Mn and Cu elements. During 850°C sintering, the alloy is mainly single γ-MnCu solid solution. At the same time, as the sintering temperature increases, the compactness and damping capacity of the alloys increases and the micro-hardness decreases.
Fig. 3 BSE images of the microstructures of the SPS730 (a), SPS770 (b), SPS810 (c) and SPS850 (d) alloys.
Mn–Cu-based alloys have received widespread attention due to their high damping properties and good mechanical properties, and have been successfully used in military and civilian fields.1–3) Its high damping capacity is mainly due to the energy dissipation of {110} twin boundaries movement. The formation of {110} twin boundaries are closely related to the antiferromagnetic transition and the phase transition from fcc cubic structure to fct tetragonal structure.4–9)
At present, Mn–Cu alloys with high Mn content are mainly obtained through casting processes. During the smelting process, the Mn element is easy to volatilize, the alloy has poor fluidity and shrinkage properties, serious component segregation, and a long material preparation cycle, which limits the practical engineering application of the material to a certain extent.
Yin et al.10) prepared Mn–Cu-based damping alloys by powder metallurgy for the first time, and found a high sub-peak at 440 K which is due to fcc-fct martensitic transformation. Luo et al.11) prepared Mn–Cu-based damping alloy by manually mixing pure metal powder, then cold pressing in a mold and then putting it into a heat treatment furnace for sintering. It is found that the alloy can form single-phase γ-MnCu only when the temperature is above 930°C. After sintering, the compactness of the alloys reached 87.2%, but there are still many α-Mn particles and incomplete sintering, which is unfavorable for the damping capacity and mechanical properties of the alloys.12) Manual mixing of pure metal powder will cause uneven mixing of ingredients. Therefore, this article attempts to use ball milling and spark plasma sintering (SPS) to prepare an alloy with a nominal composition of Mn–25Cu–2Al. The effect of adding a certain amount of Al is to improve the damping capacity and corrosion resistance of the alloy.13,14)
Ball milling can make the metal powder mix more uniformly. In the spark plasma sintering process, a certain pressure which promotes diffusion and reaction among elements is applied to quickly obtain alloys with homogeneous and fine grain size, and also higher compactness at a lower sintering temperature.15) At present, there is no relevant research on the preparation of Mn–Cu based damping alloys by SPS technology. The effect of SPS temperature on the phase composition, microstructure, compactness, microhardness, damping capacity of these materials were systematically investigated in this work.
Commercially available electrolytic Mn (99.5%, 72 g), electrolytic Cu (99.9%, 25 g), electrolytic Al (99.9%, 3 g) elemental powders (48∼150 µm) with nominal composition of Mn–25Cu–2Al were milling in a planetary ball miller (QM-QX2) for 4 h at 350 rpm under high purity argon atmosphere, stainless steel vials and balls were used as milling media, and the ball to powder mass ratio was 10:1. Stearic acid was added as process controlling agent during the milling process. The milled powders were consolidated by SPS (LABOX-350) with constant axial pressure of 30 Mpa, the cylindrical samples with the diameter of 30 mm and thickness of about 8 mm were produced by sintering for 10 min at target temperature (730, 770, 810 and 850°C, labelled as SPS730, SPS770, SPS810 and SPS850, respectively).
Compactness of sintered alloys was the ratio of density to theoretical density, and density of sintered alloys was measured based on the Archimedes method. Phase constitutions of the milled powders and the samples were identified by X-ray diffraction equipment (XRD, DX-2700B) with Cu Kα radiation (λ = 0.154598 nm). Microstructure of the milled powders and bulk samples were observed by scanning electron microscopy (SEM, Inspect F50, FEI, USA) using backscattered electron (BSE) imaging and sometimes by transmission electron microscopy (TEM, Talos F200X, FEI, USA) and high resolution TEM (HRTEM). TEM samples were prepared by mechanical polishing to less than 100 µm, followed by argon ion milling in an ion thinning instrument (Gatan 691, USA). Energy-dispersive x-ray spectroscopy (EDS, Inca X-sight) was also employed to determine the phase composition. Micro-hardness (HV) of the samples were measured by Vickers hardness tester (MHVD-50AP) with a load of 1 kg using 6 indents for each alloy at ambient temperature.
The specimens used for damping capacity testing were cut into 1.5 mm × 1.5 mm × 30 mm by electro-discharge machining (EDM). The temperature is raised from −60°C to 130°C, the heating rate is 2.5°C·min−1, the frequency is 1 Hz, and the strain amplitude is 500 × 10−6. At the same time, the variation of the damping capacity of the alloy with the strain amplitude (0∼600 × 10−6) at room temperature was also tested, with a frequency of 1 Hz.
Figure 1 show the XRD and BSE patterns of the mixed metal powder after ball milling for 4 hours. As shown in Fig. 1(a), it can be seen that the mixed metal powder after ball milling are mainly composed of α-Mn and γ-Cu. The diffraction peak of Al particles is not found in the XRD diffraction pattern, which is due to the low content of in powder mixture.16) It can be seen from Fig. 1(b) that the mixed metal powder after ball milling is significantly refined compared to pure metal powder (48∼150 µm in size). There are partially isolated irregular large-sized α-Mn (gray, 6 ± 2 µm), small size α-Mn and γ-Cu (bright, 2 ± 1 µm). It was found that the Mn content in α-Mn was 97.4% by EDS. Small size (nano-level) γ-Cu and α-Mn particles appear agglomeration phenomenon (Fig. 1(c), bottom right corner). During the ball milling process, the metal powder undergoes physical movements such as impact, extrusion, and rubbing, Mn and Cu elements diffuse each other, forming γ-Cu and α-Mn.16)
XRD pattern (a) and BSE images (b), (c) of the microstructure of the powder mixture upon ball milling of pure elemental powders: (c) shows the locally magnified microstructure of powder mixture in (b).
Figure 2 shows the XRD pattern of the sintered alloys. It can be seen from Fig. 2(a) that the SPS730 alloy consists of α-Mn and γ-MnCu phases. With the increase of the sintering temperature, the diffraction peak intensity of α-Mn phase decreased in the SPS770 alloy. In the SPS810 alloy, the diffraction peak of α-Mn phase has disappeared, and only the diffraction peak of single-phase γ-MnCu exists. When the sintering temperature is 850°C, the diffraction peak of the (220) crystal plane split, forming two special diffraction peaks of (220) and (202) (Fig. 2(b)). At the same time, it is found that the diffraction peak of the (111) crystal plane tends to shift to a high angle direction as the sintering temperature increases. Indicating that the degree of solid solution of the alloy is not complete during low sintering temperature. With the increase of sintering temperature, the diffusion of Al and especially Mn into γ-Cu lattices, the degree of solid solution increases. The existing single-phase γ-MnCu in SPS810 alloy indicates that the alloy has been completely dissolved. In the SPS850 alloy, the (220) and (202) diffraction peaks formed by the splitting of the characteristic diffraction peaks of the (220) crystal plane indicate that the fcc to fct structure transformation has occurred, and the martensite structure of the fct structure is formed.12,17,18)
XRD patterns of the SPS alloys (a); (b) shows the locally magnified XRD patterns of the SPS850 alloy.
Table 1 shows the compactness of sintered alloys. As the sintering temperature increases, the compactness increases. The compactness of the SPS730 alloy is 95.6%, and the compactness of the SPS850 alloy increases to 98.7%, significantly better than the 87.2% mentioned earlier.12) SPS provides sufficient activation energy for the surface diffusion of the powder particles to the volume diffusion, which can reduce or even eliminate the connection holes between the particles, and increase the sintering compactness.
Microstructures of the sintered alloys are shown in Fig. 3. The SPS730 alloy is mainly composed of γ-MnCu (gray phase, matrix structure, 92.9Mn–3.7Cu–3.4Al, at%, Table 2), γ-CuMn (bright phase, mainly in strip shape, 11.6Mn–87.8Cu–0.6Al, at%, Table 2) and α-Mn (dark gray, relatively concentrated). With the increase of sintering temperature, γ-CuMn and α-Mn gradually disappeared. In SPS770 alloy, the size of γ-CuMn phase is obviously reduced, the number decreases, the proportion of α-Mn phase becomes shallower, and there is a tendency to disperse. In the SPS810 alloy, the proportion of α-Mn phase is further reduced, and the tendency of dispersion is more obvious. In SPS850 alloys, α-Mn and γ-CuMn phases are rarely found, and the alloy is mainly γ-MnCu (75.1Mn–24.8Cu–1.1Al, at%, Table 2). At the same time, the high-magnification BSE images (located above the picture) show that when sintered at 730°C, it appears as an irregular combination of solid particles. As the sintering temperature increases, the sharp corners between the irregular particles melt, and the crystal grains become more regular round. When sintered at 850°C, pores formed by sintering appear between the crystal grains, and the crystal grains have a tendency to grow. Combining XRD, BSE and EDS of different SPS alloys (Table 2), the increase of sintering temperature promotes the diffusion of metal elements and particles melting, which causes the Mn element in α-Mn to diffuse to the surroundings, and the γ-CuMn phase obtains more Mn from nearby, which makes the composition of the alloy tend to be uniform, and finally forms a single-phase γ-MnCu matrix.
BSE images of the microstructures of the SPS730 (a), SPS770 (b), SPS810 (c) and SPS850 (d) alloys.
Table 3 shows the micro-hardness of sintered alloys. It can be seen that the hardness of SPS730 alloy is the highest. As the sintering temperature increases, the hardness of the alloy decreases. SPS730 alloy is composed of γ-MnCu, γ-CuMn and α-Mn, among which α-Mn has a high hardness value.19,20) At the same time, the grain size of the alloy is small and the dislocation defects formed and micro-strain in the ball milling did not completely disappear, and there was a fault strengthening effect,16) so the hardness value of SPS730 alloy is highest. As the temperature rises, Mn and Cu elements diffuse each other, and the composition tends to be uniform, forming a single-phase γ-MnCu. At the same time, due to the growth of crystal grains, SPS850 alloy has the lowest hardness value.
Figure 4(a) shows the relationship between the damping capacity of the SPS alloys with the torsional strain at room temperature and a frequency of 1 Hz. It can be seen that the damping capacity of the alloys increases as the strain increases. And with the increase of the sintering temperature, the damping capacity of the higher temperature sintered alloys is better than the lower temperature sintered alloys in the entire strain range. Figure 4(b) shows the variation of the damping capacity of the SPS alloys with temperature at a frequency of 1 Hz and an applied strain amplitude of 5 × 10−4. It can be seen that in the low temperature range (−60°C to 50°C), the damping capacity of the higher temperature sintered alloys is better than the lower temperature sintered alloys. And a typical twin relaxation peak appears in the low temperature section, and a martensitic transformation peak appears in the high temperature section.21–23) It is found that the phase transition peak of SPS730 alloy is at 46.6°C, SPS770 alloy is 62.1°C, and SPS810 alloy is 118.3°C. It can be inferred that the phase transition peak of SPS850 alloy is at a higher temperature. Yin et al. also found a high sub-peak at 440 K which is due to fcc-fct martensitic transformation.10)
(a) Variation of the damping capacity of the SPS alloys with applied strain amplitude at room temperature and a frequency of 1 Hz; (b) Variation of the damping capacity of the SPS alloys with temperature at a frequency of 1 Hz and an applied strain amplitude of 5 × 10−4.
The damping capacity of Mn–Cu-based damping alloys is due to the energy dissipation of twin boundaries induced by antiferromagnetic transformation and fcc-fct phase transformation.24–26) There are many γ-CuMn and α-Mn grains in the SPS730 alloy. The formation of {101} twins is closely related to martensitic transformation, and the low content of Mn in γ-CuMn leads to the decrease of martensitic transformation temperature. Meanwhile, the Cu-rich phase in γ-CuMn will become the resistance of martensitic transformation in the Mn-rich region, and delay the formation of martensitic twins.27) The production of α-Mn reduces the number of Mn-rich regions, leads to a decrease in martensite transition temperature, and impedes twin formation.26,28) Furthermore, the movement the {101} twin boundaries is restricted due to existence of visible α-Mn grains in the SPS730 alloy. However, which has almost not occurred in the SPS850 alloy, so the damping capacity of SPS730 alloy is the lowest in this study. It was found that the typical tweed structure and the fine twins from TEM image of the SPS850 alloy, as shown in Fig. 5(a) and Fig. 5(b). The tweed structure is the early stage of twin formation without obvious interface.29,30) It is the amplitude modulation decomposition structure, which promotes the formation of Mn-rich regions and provides advantage conditions for martensitic transformation. The high-resolution TEM image (HRTEM) in Fig. 5(c) and corresponding IFFTs images and SAED pattern of the Fig. 5(c) region (Fig. 5(d)). The results suggest that the atoms are symmetrically distributed along the twinning plane, confirming the existence of fine twins.
Typical TEM image of mottled substructure (a) and twin structures (b) in the SPS850 alloy, induced by fcc-fct phase transformation; HREM image of the SPS850 alloy (c) and corresponding IFFTs images and SAED pattern (d) of the region in (c) clearly showing the Atoms are symmetrically distributed along the twinning plane.
During the cooling of the alloy, the Mn-rich regions undergoes a structural transformation from face-centered cubic to tetragonal. The relationship between the distortion degree c/a and the manganese content cMn satisfies the following formula:18,31–33)
| \begin{equation} c/a = 2.638 - 3.317c_{Mn} + 1.618c_{\textit{Mn}}^{2} \end{equation} | (1) |
| \begin{equation} 1/d^{2} = h^{2} + k^{2}/a^{2} + l^{2}/c^{2} \end{equation} | (2) |
| \begin{equation} 2d\sin \theta = \lambda \end{equation} | (3) |
According to the XRD diffraction pattern of Fig. 2(b), the relationship between the Mn content and the lattice constant in the Mn-rich regions of the SPS850 alloy can be obtained, as shown in Table 4. It can be seen that the Mn content in the Mn-rich zone reaches 85.74%, which exceeds the critical Mn content for fcc-fct transformation by 83.4%. Therefore, this type of Mn-rich zone undergoes fcc-fct structural transformation at room temperature, forming martensite twins and obtaining damping capacity.
In the present study, a Mn–Cu-based damping alloy with a nominal composition of Mn–25Cu–2Al was produced by ball milling and SPS technology. The phase composition, microstructure, compactness, micro-hardness, damping capacity of these materials were investigated. The main conclusions obtained are as follows:
This study was financially supported by Opening Foundation of Sichuan Province Engineering Research Center for Powder Metallurgy, Chengdu University.
