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A comparative study on microstructure and damping capacity of Mn-Cu based alloys with dendrite and equiaxial grain
Vacuum 168 (2019) 108814
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A comparative study on microstructure and damping capacity of Mn-Cu based alloys with dendrite and equiaxial grain
T
Song Zhanga,*, Xiping Guob, Ye Tangb, Shuai Zhonga, Yonggang Xua a
Key Laboratory of Advanced Technologies of Materials, Ministry of Education, School of Materials Science and Engineering, Southwest Jiaotong University, Chengdu, 610031, PR China b State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi'an, 710072, PR China
A R T I C LE I N FO
A B S T R A C T
Keywords: Mn-Cu based alloy Damping capacity Phase transformation damping Twin relaxation Internal friction Dendritic segregation
A 70Mn-24.95Cu-3Al-2Zn-0.05Ce (at.%) alloy was prepared by vacuum induction melting (VIM) technology. Part of the alloy was further hot rolled and annealed (HRA) respectively at 500 °C with a thickness reduction of 50% and 850 °C for 5 h. Then, VIM and HRA samples were aged at 430 °C for 2–4 h. The comparative study on microstructure and damping capacity of VIM and HRA alloys has been conducted. The results show that the microstructure of VIM alloy comprises γ-MnCu dendrites with obvious compositional segregation while that of HRA alloy primarily consists of equiaxial γ-MnCu grains. Some α-Mn particles are formed in HRA alloy aged for 4 h. The dendritic segregation has no obvious effect on starting fcc-fct transformation temperature of the alloys. The VIM alloy has larger tetragonal distortion 1-c/a than HRA one. Two stepped fcc-fct phase transformation occurs and couples to promote the formation of a high and wide phase transformation damping platform in VIM alloy especially aged for 4 h compared to a phase transformation damping peak in HRA one. There is a significantly towering twin relaxation internal frication peak in VIM alloy especially aged for 2 h compared to HRA one.
1. Introduction As a typical class of functional and structural materials, Mn-Cu based alloys have attracted much attention and possessed broad application prospects in many fields such as automobile, rail transit and military equipment etc. because of their high damping capacity and favorable mechanical properties [1]. The damping capacity of Mn-Cu based alloys has mainly been achieved by the movement and relaxation of {101} twin boundaries induced by martensite phase transformation of fcc γ-MnCu to fct γ′-MnCu [2]. Generally, the fcc-fct phase transformation temperature which is linearly dependent on Mn content in the alloy can be raised by aging treatment due to the formation of spinodal decomposition microstructure composed of Mn-rich and Curich regions with sizes of 101 nm order, thus promoting the formation of abundant {101} twins and high damping capacity above room temperature [3,4]. Mn-Cu based damping alloys have usually been prepared by vacuum induction melting, subsequent plastic working (hot-forging or hotrolling) and further heat-treatment (annealing and aging) in order to obtain homogeneous equiaxial grains [5]. Recently, Zhong et al. [6], Liu et al. [7] and Yin et al. [8] have indicated that the direct vacuum
*
induction melted Mn-Cu based alloys still possess excellent damping capacity which is even higher than that of the plastically worked and annealed ones upon appropriate aging treatment. However, they have not presented an exact interpretation for this result. It is well known that the microstructure of the plastically worked and annealed alloys is generally composed of homogeneous equiaxial grains while that of the vacuum induction melted ones comprises dendrites with obvious compositional segregation, and also the dendritic and interdendritic zones own relatively high and low Mn concentrations respectively [6–8]. Considering the dependence of damping capacity on Mn content [3,4], weather the high damping characteristic of the vacuum induction melted alloys is closely related to their dendritic segregations? If so, what's the difference of damping behavior between plastically worked and annealed and vacuum induction melted alloys, such as in terms of variation of damping capacity with temperature or strain amplitude, and coupling characteristic of fcc-fct phase transformation damping and twin relaxation internal friction and so on? All of these have not been understood clearly. Therefore, further study is need. According to the open literatures, the addition of Al improves the casting performance (such as flowability and contractility of the alloy
Corresponding author. E-mail address: [email protected] (S. Zhang).
https://doi.org/10.1016/j.vacuum.2019.108814 Received 23 April 2019; Received in revised form 3 June 2019; Accepted 9 July 2019 Available online 10 July 2019 0042-207X/ © 2019 Elsevier Ltd. All rights reserved.
Vacuum 168 (2019) 108814
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Fig. 1. Typical BSE images of the microstructures of Mn-Cu based alloys: (a) VIM2h and (b) HRA2h.
2000#, polished and cleaned in an ultrasonic alcohol bath, and eventually blown-dry successively. The constituent phases of the samples were identified by X-ray diffraction analysis using ICDD cards (XRD, Panalytical X'Pert PRO, Cu Kα, scanning rate 0.05°/s). Also, the XRD results were used to calculate lattice parameters of constituent phases by analyzing their diffraction angles. The microstructure of the samples was observed by scanning electron microscopy (SEM, QUANTA FEG 250) using backscattered electron (BSE) imaging, and sometimes by transformation electron microscopy (TEM, Tecnai G2 F30 operating at 300 kV). Thin foils for TEM observation were fabricated by a conventional method, i.e. the disks with a thickness of 300 μm were cut from samples by EDM, then mechanically ground down to 50 μm and dimpled to 10 μm, and finally ion-milled. Energy dispersive X-ray spectroscopy (EDS, Inca X-sight) was employed to determine the phase composition. In addition, the quantitative analysis of the constituent phases was carried out using Image-Pro Plus 6.0 high resolution imaging analysis software. At least 3 BSE images at the same magnification (200 × ) were used for each sample condition. The damping capacity (Q−1, tangent value of phase angle difference between stress and strain) and relative elastic modulus (Er, corresponding to Young's modulus one by one) of the samples with a size of 1.5 × 1.5 × 50 mm3 by EDM were examined in an internal friction instrument using inverted torsion pendulum mode. The measurement had been carried out during heating from −80 to 200 °C at a temperature raise rate of 2.5 °C/min, and with a frequency of 1 Hz and a strain amplitude of 5 × 10−4. It should be pointed out that the curve of variation of relative elastic modulus with temperature, i.e. Er-T curve can be used to assess the fcc-fct transformation temperature of the alloys studied in this work [14]. Besides, the variation of damping capacity of the samples with strain amplitude (up to the order of 10−3) was also obtained at room temperature with a frequency of 1 Hz.
melt), corrosion resistance and mechanical properties of the alloys, but degrades their damping capacity to some extent [9,10]. The element Zn is considered to be beneficial to both ameliorate the damping capacity and increase the strength of Mn-Cu alloys [11,12]. The Ce microalloying can suppress the enrichment of impurity elements (such as C and S etc.) and the formation of α-Mn particles at the grain and twin boundaries, thus improving the mobility of these boundaries and damping capacity of the alloys [2]. Therefore, an optimized Mn-Cu-AlZn-Ce system alloy was selected, and also the vacuum induction melting technology, hot rolling and annealing processing were employed to prepare the alloys with dendrites and homogeneous equiaxial grains in the present work. The microstructure and damping capacity of Mn-Cu based alloys with dendrites and equiaxial grains, especially the effects of compositional segregation on damping behavior of the alloys have been investigated and evaluated. For the convenience of description, the vacuum induction melted as well as hot rolled and annealed alloys are referred to VIM and HRA ones hereinafter, respectively. This paper will contribute to guiding the design and development of novel as-cast Mn-Cu based alloys, especially some complicated casting components used for controlling vibration and reducing noise. 2. Experimental procedures The ingot of Mn-Cu based alloy with a nominal composition of 70Mn-24.95Cu-3Al-2Zn-0.05Ce (at.%) was prepared by vacuum induction melting technology under an argon atmosphere with a purity of 99.999 vol%. Notably, the pressure of vacuum chamber had been pumped below 0.1 Pa before filling with argon, otherwise probably resulting in the formation of Mn oxides and thus degrading the damping capacity and mechanical properties of the alloys [4,13]. The alloy was melted in a magnesia crucible and then poured into a metal mold for cooling. Many rods with a size of 10 × 10 × 100 mm3 were cut at the same distance from the center axis of the cylindrical alloy ingot with a size of about Φ100 × 120 mm3 by electro-discharge machining (EDM). Some of alloy rods were put in a box-type resistance furnace and heated up at a temperature raise rate of 5 °C/min. When the temperature reached 500 °C, they were immediately taken out and hot rolled into strips on a rolling mill with a thickness reduction of 50%. Further, the hot rolled alloy strips were annealed at 850 °C for 5 h in a box-type resistance furnace and then water quenched. To obtain the spinodal decomposition microstructure and increase the fcc-fct transformation temperatures, the VIM rods and HRA strips were both aged at 430 °C for 2 and 4 h in a box-type resistance furnace and then water quenched. The selection of aging temperature and time is primarily based on the previous experimental results from our research group. In this work, the Mn-Cu based alloys aged for different time are referred to VIM2h, VIM4h, HRA2h and HRA4h ones hereinafter, respectively. The samples used for microstructural analysis were cut into 8 × 8 × 5 mm3 square blocks from the aged VIM rods and HRA strips by EDM. Then, they were ground using SiC-grit abrasive papers down to
3. Results and discussion 3.1. Microstructure of VIM and HRA alloys Figs. 1 and 2 show the BSE images and XRD patterns of the Mn-Cu based alloys upon aging treatment, respectively. It can be seen that the microstructure of the VIM alloy mainly comprises γ-MnCu dendrites with distinct compositional segregation between dendritic and interdendritic, i.e. dark and light zones with area fractions of about 73% and 27% respectively (Fig. 1a). The EDS analysis results show that the compositions of dendritic and interdendritic zones are 76.9Mn-18.6Cu3.1Al-1.4Zn and 61.3Mn-34.4Cu-2.8Al-1.5Zn (at.%) respectively, indicating the primarily Mn and Cu segregations. Furthermore, Mn is evidently predominant in the initially solidified dendritic zone compared to terminally solidified interdendritic one due to the higher melting temperature of Mn than Cu. In contrast, the microstructure of the HRA alloy is mainly composed of homogeneous equiaxial γ-MnCu grains with an approximate composition of 71.2Mn-24.3Cu-2.9Al-1.6Zn 2
Vacuum 168 (2019) 108814
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the alloys. In this work, both the (220) and (202) peaks have been fitted using quadratic polynomial function, i.e. I = M(2θ)2 + N(2θ) + L, and the corresponding coefficients and parameters calculated by XRD analysis are presented in Table 1. It can be found that the 1-c/a value of VIM alloy is higher than that of HRA one. Moreover, with increasing the aging time from 2 to 4 h, the 1-c/a value of the former alloy increases while that of the latter one deceases instead (0.0246, 0.0303, 0.0208 and 0.0066 for VIM2h, VIM4h, HRA2h and HRA4h alloys respectively). Generally, the higher 1-c/a value, i.e. larger tetragonal distortion implies more {101} twins induced by fcc-fct transformation and thus higher damping capacity of the alloy [15]. As mentioned above, the αMn phase is inclined to form in the HRA alloy especially after aging for 4 h (Figs. 2 and 3). Also considering that the HRA4h alloy has the lowest 1-c/a value, it is easily concluded that the existence of α-Mn phase certainly slows down the fcc-fct phase transformation and therefore the formation of {101} twins in this alloy to some extent [2,4].
Fig. 2. XRD patterns of Mn-Cu based alloys.
(at.%), as shown in Fig. 1b. It should be noted that the element Ce has not been determined by EDS analysis, which is mainly associated with its very small amount in the VIM and HRA alloys. The several diffraction peaks of α-Mn phase are detected only in the XRD pattern of the HRA4h alloy (Fig. 2). It is implied that the overaging which means the formation of α-Mn with increasing the aging time, prefers to occur in the plastically worked alloy instead of vacuum induction melted one. The fine α-Mn particles with sizes in nanoscale are distributed within γ-MnCu grains and also at γ-MnCu grain boundaries, as shown in Fig. 3. Furthermore, some high density dislocations generating during hot rolling are still observed around these α-Mn particles (as indicated in Fig. 3). It is therefore inferred that the distortion regions caused by high density dislocations would have provided nuclei for the formation of α-Mn phase in HRA alloy during aging. Fig. 4 shows the locally magnified XRD patterns of VIM and HRA alloys. It can be seen that the original (220) diffraction peak separates into both (220) and (202) ones in the XRD patterns of the two alloys irrespective of aging time, which means that the fcc-fct phase transformation occurs and {101} twins form [15]. The typical twinning bands with different orientations from VIM2h alloy are displayed in Fig. 5. Moreover, there is an intersection angle of about 60° between twinning bands which is very close to the intersection angle between {101} boundaries in different fct twin variants [16], calculated approximately by the following equation usually applied to fcc structure (fct and fcc structures have very small difference of lattice parameters):
cos θ =
3.2. Damping capacity of VIM and HRA alloys Fig. 6 shows the variation of the damping capacity and relative elastic modulus of VIM and HRA alloys with temperature at a strain amplitude of 5 × 10−4. It can be seen from Q−1-T curves that the VIM alloy possesses significantly higher damping capacity than HRA one irrespective of aging time. Generally, there is a phase transformation damping peak and a twin relaxation internal friction one at high and low temperature sides in the Q−1-T curves of Mn-Cu based alloys, respectively [17]. In this work, a nontypical phase transformation peak, namely a phase transformation damping platform occurs at high temperature side in the Q−1-T curves of the two VIM alloys (Fig. 6a). In contrast, there is a relatively typical phase transformation peak in the Q−1-T curve of HRA2h alloy, however which is not observed in that of HRA4h one again due to its very low phase transformation damping capacity (Fig. 6a). As mentioned in section 3.1, many fine α-Mn particles formed in HRA4h alloy slow down the fcc-fct phase transformation and the formation of {101} twins, thus leading to the lowest damping capacity of this alloy. In addition, the obviously towering twin relaxation peaks are displayed in the Q−1-T curves of VIM alloys compared to HRA ones (Fig. 6a). It should be noted that the phase transformation damping capacity of VIM2h alloy is lower than that of VIM4h one, but the twin relaxation internal frication of the former alloy is evidently higher than that of the latter one (Fig. 6a). Fig. 7 shows the variation of damping capacity of VIM and HRA alloys with strain amplitude at room temperature. It is also observed that the VIM alloys have higher damping capacity than HRA alloys over the whole strain amplitude range irrespective of aging time. Moreover, the damping capacity of both alloys decreases with increasing the aging time from 2 to 4 h to some extent. It is well known that the temperature corresponding to minimum elastic modulus in E-T curve represents the starting phase transformation one of Mn-Cu based alloys [18]. As observed in Fig. 6b, there are similar modulus valley temperatures, i.e. starting fcc-fct transformation ones in Er-T curves of VIM and HRA alloys and they raise with increase in aging time (116.1 vs 122.6 and 146.2 vs 155.0 °C for VIM2h vs HRA2h and VIM4h vs HRA4h alloys, respectively). The similar starting fcc-fct transformation temperatures are probably related to the relatively close average Mn concentrations of Mn-rich regions in the decomposition microstructures from dendritic zone for VIM alloy and equiaxial grain for HRA one, respectively. The Mn concentration of Mnrich regions can be approximately counted as 88.8 vs 89.1 and 91.0 vs 91.7 at.% for VIM2h vs HRA2h and VIM4h vs HRA4h alloys respectively by the following empirical equation [3,4]:
H1 H2 + K1 K2 + L1 L2 H12
+ K12 + L12 ⋅ H22 + K22 + L22
(1)
where θ and (HKL) are crystal plane intersection angle and crystal plane indices respectively. Notably, the separation extent of (220) diffraction peak in the XRD pattern differs from each other (Fig. 4), indicating the various axial ratios (c/a) and tetragonal distortions (1-c/a) of fct γ′-MnCu phase in
Fig. 3. Typical TEM images from HRA4h alloy, showing the fine α-Mn particles distributed within γ-MnCu grains (a) and also at γ-MnCu grain boundaries (b), and high density dislocations around these particles as well.
Ms = 1264.8cMn – 731.8
(2)
where Ms represents starting phase transformation temperature (K) and 3
Vacuum 168 (2019) 108814
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Fig. 4. Locally magnified XRD patterns of Mn-Cu based alloys, showing the separation of (220) diffraction peak into (220) and (202) ones: (a) VIM2h, (b) VIM4h, (c) HRA2h and (d) HRA4h.
cMn is Mn concentration (at.%). It should be pointed out that the fcc-fct phase transformation occurs firstly in Mn-rich region in the decomposition microstructure of dark dendritc zone and then in that of light interdendritic one, since the former zone has higher Mn concentration than the latter one in the VIM alloy (76.9 and 61.3 at.% respectively) [3,4], as shown in Fig. 1a. In other words, the two stepped fcc-fct phase transformation happens and couples to each other in the VIM alloy with obvious compositional segregation, which inevitably imposes an important influence on its damping capacity, especially phase transformation damping capacity. This illustrates the phase transformation damping platform in the Q−1T curve of the VIM alloy in comparison with the phase transformation damping peak of the HRA alloy (Fig. 6a). Furthermore, with increasing the aging time from 2 to 4 h, the starting phase transformation temperature increases due to the more sufficient spinodal decomposition [5,15], thus resulting in a higher and wider phase transformation damping platform in VIM4h alloy than in VIM2h one (Fig. 6a). On one hand, the high fcc-fct transformation temperature promotes the formation of fct {101} twins, which improves the damping capacity of Mn-Cu based alloys [19]. On the other hand, however, the axial ratio c/a of fct phase goes down, namely the tetragonal distortion (1-c/a) increases with decrease in temperature so that the {101} planes
Fig. 5. TEM image of typical twinning bands from VIM2h alloy, showing an intersection angle of about 60° between twinning bands with different orientations.
Table 1 The coefficients M, N, L in quadratic polynomial function, i.e. I = M(2θ)2 + N(2θ) + L fitted for separated (220) and (202) peaks and corresponding peak 2θ in Fig. 4, and the lattice parameters c and a, axial ratio c/a and tetragonal distortion 1-c/a of fct γ′-MnCu phase in the alloys calculated by XRD analysis. Alloy
Peak
M
N
L
Peak 2θ
c
a
c/a
1-c/a
VIM2h
(220) (202) (220) (202) (220) (202) (220) (202)
−1415.9 −493.4 −1303.8 −1305.1 −224.7 −1016.6 −7165.9 −2050.2
200537.3 70906.4 184486.5 188030.3 31966.4 146410.5 1012020.0 290620.3
−7.1E6 −2.5E6 −6.5E6 −6.8E6 −1.1E6 −5.3E6 −3.6E7 −1.0E7
70.8162 71.8549 70.7495 72.0367 71.1313 72.0099 70.6136 70.8830
3.6707
3.7634
0.9754
0.0246
3.6522
3.7665
0.9697
0.0303
3.6708
3.7489
0.9792
0.0208
3.7480
3.7728
0.9934
0.0066
VIM4h HRA2h HRA4h
4
Vacuum 168 (2019) 108814
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Fig. 6. Variation of damping capacity (a) and relative modulus (b) of Mn-Cu based alloys with temperature at a strain amplitude of 5 × 10−4.
4. Conclusions (1) The microstructure of VIM alloy mainly consists of γ-MnCu dendrites with obvious compositional segregation while that of HRA alloy primarily comprises equiaxial γ-MnCu grains. Also, some very fine α-Mn particles are formed in HRA alloy aged for 4 h. (2) The dendrtic segregation has no significant influence on the starting fcc-fct transformation temperature of Mn-Cu based alloys which is primarily related to aging time. However, the VIM alloy has larger tetragonal distortion 1-c/a than HRA one. Moreover, accompanying the raise in aging time, 1-c/a value of the former alloy increases while that of the latter one deceases. (3) The two stepped fcc-fct transformation occurs and couples to promote the formation of a phase transformation damping platform in VIM alloy compared to a phase transformation damping peak in HRA one, which is still enhanced with increase in aging time. (4) A distinctly towering twin relaxation internal frication peak is present in VIM alloy especially aged for 2 h because of the synergetic damping effect between {101} twins in the dendritic and interdendritic zones, compared to HRA one.
Fig. 7. Variation of damping capacity of Mn-Cu based alloys with strain amplitude at room temperature.
respectively in the two twinned constituent grains might become separated from each other and incoherent. In order to maintain the coherent {101} twin boundary, the initially parallel {101} planes in these two constituent grains then rotate in opposite directions. As thus, the initial orientation relationship between the two twinned constituent grains is destroyed, i.e. their relative reorientation occurs, which results in the reduction of damping capacity at lower temperatures while the twin boundary relaxation is explained by a thermal activation process [19]. In this work, owing to having higher starting fcc-fct transformation temperature, the tetragonal distortion (1-c/a) of the VIM4h alloy is larger than that of the VIM2h one (0.0303 and 0.0246 respectively, Table 1). Therefore, a bigger destruction of initial twinning relationship and a more evident relative reorientation between the two fct twinned constituent grains are present in the former alloy than in the latter one. This illustrates the obviously lower twin relaxation internal friction of VIM4h alloy than VIM2h one to some extent (Fig. 6a). As mentioned above, the starting phase transformation temperature of Mn-Cu based alloys is not significantly influenced by the compositional segregation (116.1 and 122.6 °C for VIM2h and HRA2h alloys with γ-MnCu dendrites and equiaxial γ-MnCu grains, respectively). Nevertheless, their damping capacity, especially twin relaxation internal friction is closely associated with such segregation (Figs. 6a and 7). Considering that the two stepped fcc-fct phase transformation occurs respectively in the dendritic and interdendritic zones of VIM alloy and the corresponding collaborative phase transformation damping exists (Fig. 6a), so it is easily imagined that the resulting {101} twins in these two zones would also have cooperated to improve its damping capacity. That is to say the synergetic effect of {101} twins in the two zones is probably responsible for higher twin relaxation internal friction of VIM2h alloy than HRA2h one. In addition, the relatively high density dislocations in some localized regions (the typical TEM image of some high density dislocations from HRA alloy is shown in Fig. 3) may also be another reason for degrading the damping capacity of HRA2h alloy.
Acknowledgments This work was supported by the National Natural Science Foundation of China (No. 51701167), Fundamental Research Funds for the Central Universities (No. 2682017CX073) and Fund of the State Key Laboratory of Solidification Processing in NWPU (No. SKLSP201825). Appendix A. Supplementary data Supplementary data to this article can be found online at https:// doi.org/10.1016/j.vacuum.2019.108814. References [1] D. Weng, S. Liu, J. He, Research and industrialization status of Mn-based damping alloys, Sci. Technol. Rev. 32 (3) (2014) 77–83. [2] F.S. Lu, B. Wu, J.F. Zhang, P. Li, D.L. Zhao, Microstructure and damping properties of MnCuNiFeCe alloy, Rare Met. 35 (8) (2016) 615–619. [3] F. Yin, Y. Ohsawa, A. Sato, K. Kawahara, Phase decomposition of the γ phase in a Mn-30 at.% Cu alloy during aging, Acta Mater. 48 (2000) 1273–1282. [4] T. Sakaguchi, F. Yin, Holding temperature dependent variation of damping capacity in a MnCuNiFe damping alloy, Scr. Mater. 54 (2006) 241–246. [5] J. Yan, N. Li, X. Fu, Y. Zhang, The strengthening effect of spinodal decomposition and twinning structure in MnCu-based alloy, Mater. Sci. Eng. A 618 (2014) 205–209. [6] Z. Zhong, W. Liu, N. Li, J. Yan, J. Xie, D. Li, Y. Liu, X. Zhao, S. Shi, Mn segregation dependence of damping capacity of as-cast M2052 alloy, Mater. Sci. Eng. A 660 (2016) 97–101. [7] W. Liu, N. Li, Z. Zhong, J. Yan, D. Li, Y. Liu, X. Zhao, S. Shi, Novel cast-aged MnCuNiFeZnAl alloy with good damping capacity and high usage temperature toward engineering application, Mater. Des. 106 (2016) 45–50. [8] F.X. Yin, S. Iwasaki, T. Sakaguchi, K. Nagai, The improved damping behavior of MnCu high damping alloy obtained by solidification process control, Prog. Phys. 26 (3–4) (2006) 323–331.
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A comparative study on microstructure and damping capacity of Mn-Cu based alloys with dendrite and equiaxial grain
T
Song Zhanga,*, Xiping Guob, Ye Tangb, Shuai Zhonga, Yonggang Xua a
Key Laboratory of Advanced Technologies of Materials, Ministry of Education, School of Materials Science and Engineering, Southwest Jiaotong University, Chengdu, 610031, PR China b State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi'an, 710072, PR China
A R T I C LE I N FO
A B S T R A C T
Keywords: Mn-Cu based alloy Damping capacity Phase transformation damping Twin relaxation Internal friction Dendritic segregation
A 70Mn-24.95Cu-3Al-2Zn-0.05Ce (at.%) alloy was prepared by vacuum induction melting (VIM) technology. Part of the alloy was further hot rolled and annealed (HRA) respectively at 500 °C with a thickness reduction of 50% and 850 °C for 5 h. Then, VIM and HRA samples were aged at 430 °C for 2–4 h. The comparative study on microstructure and damping capacity of VIM and HRA alloys has been conducted. The results show that the microstructure of VIM alloy comprises γ-MnCu dendrites with obvious compositional segregation while that of HRA alloy primarily consists of equiaxial γ-MnCu grains. Some α-Mn particles are formed in HRA alloy aged for 4 h. The dendritic segregation has no obvious effect on starting fcc-fct transformation temperature of the alloys. The VIM alloy has larger tetragonal distortion 1-c/a than HRA one. Two stepped fcc-fct phase transformation occurs and couples to promote the formation of a high and wide phase transformation damping platform in VIM alloy especially aged for 4 h compared to a phase transformation damping peak in HRA one. There is a significantly towering twin relaxation internal frication peak in VIM alloy especially aged for 2 h compared to HRA one.
1. Introduction As a typical class of functional and structural materials, Mn-Cu based alloys have attracted much attention and possessed broad application prospects in many fields such as automobile, rail transit and military equipment etc. because of their high damping capacity and favorable mechanical properties [1]. The damping capacity of Mn-Cu based alloys has mainly been achieved by the movement and relaxation of {101} twin boundaries induced by martensite phase transformation of fcc γ-MnCu to fct γ′-MnCu [2]. Generally, the fcc-fct phase transformation temperature which is linearly dependent on Mn content in the alloy can be raised by aging treatment due to the formation of spinodal decomposition microstructure composed of Mn-rich and Curich regions with sizes of 101 nm order, thus promoting the formation of abundant {101} twins and high damping capacity above room temperature [3,4]. Mn-Cu based damping alloys have usually been prepared by vacuum induction melting, subsequent plastic working (hot-forging or hotrolling) and further heat-treatment (annealing and aging) in order to obtain homogeneous equiaxial grains [5]. Recently, Zhong et al. [6], Liu et al. [7] and Yin et al. [8] have indicated that the direct vacuum
*
induction melted Mn-Cu based alloys still possess excellent damping capacity which is even higher than that of the plastically worked and annealed ones upon appropriate aging treatment. However, they have not presented an exact interpretation for this result. It is well known that the microstructure of the plastically worked and annealed alloys is generally composed of homogeneous equiaxial grains while that of the vacuum induction melted ones comprises dendrites with obvious compositional segregation, and also the dendritic and interdendritic zones own relatively high and low Mn concentrations respectively [6–8]. Considering the dependence of damping capacity on Mn content [3,4], weather the high damping characteristic of the vacuum induction melted alloys is closely related to their dendritic segregations? If so, what's the difference of damping behavior between plastically worked and annealed and vacuum induction melted alloys, such as in terms of variation of damping capacity with temperature or strain amplitude, and coupling characteristic of fcc-fct phase transformation damping and twin relaxation internal friction and so on? All of these have not been understood clearly. Therefore, further study is need. According to the open literatures, the addition of Al improves the casting performance (such as flowability and contractility of the alloy
Corresponding author. E-mail address: [email protected] (S. Zhang).
https://doi.org/10.1016/j.vacuum.2019.108814 Received 23 April 2019; Received in revised form 3 June 2019; Accepted 9 July 2019 Available online 10 July 2019 0042-207X/ © 2019 Elsevier Ltd. All rights reserved.
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Fig. 1. Typical BSE images of the microstructures of Mn-Cu based alloys: (a) VIM2h and (b) HRA2h.
2000#, polished and cleaned in an ultrasonic alcohol bath, and eventually blown-dry successively. The constituent phases of the samples were identified by X-ray diffraction analysis using ICDD cards (XRD, Panalytical X'Pert PRO, Cu Kα, scanning rate 0.05°/s). Also, the XRD results were used to calculate lattice parameters of constituent phases by analyzing their diffraction angles. The microstructure of the samples was observed by scanning electron microscopy (SEM, QUANTA FEG 250) using backscattered electron (BSE) imaging, and sometimes by transformation electron microscopy (TEM, Tecnai G2 F30 operating at 300 kV). Thin foils for TEM observation were fabricated by a conventional method, i.e. the disks with a thickness of 300 μm were cut from samples by EDM, then mechanically ground down to 50 μm and dimpled to 10 μm, and finally ion-milled. Energy dispersive X-ray spectroscopy (EDS, Inca X-sight) was employed to determine the phase composition. In addition, the quantitative analysis of the constituent phases was carried out using Image-Pro Plus 6.0 high resolution imaging analysis software. At least 3 BSE images at the same magnification (200 × ) were used for each sample condition. The damping capacity (Q−1, tangent value of phase angle difference between stress and strain) and relative elastic modulus (Er, corresponding to Young's modulus one by one) of the samples with a size of 1.5 × 1.5 × 50 mm3 by EDM were examined in an internal friction instrument using inverted torsion pendulum mode. The measurement had been carried out during heating from −80 to 200 °C at a temperature raise rate of 2.5 °C/min, and with a frequency of 1 Hz and a strain amplitude of 5 × 10−4. It should be pointed out that the curve of variation of relative elastic modulus with temperature, i.e. Er-T curve can be used to assess the fcc-fct transformation temperature of the alloys studied in this work [14]. Besides, the variation of damping capacity of the samples with strain amplitude (up to the order of 10−3) was also obtained at room temperature with a frequency of 1 Hz.
melt), corrosion resistance and mechanical properties of the alloys, but degrades their damping capacity to some extent [9,10]. The element Zn is considered to be beneficial to both ameliorate the damping capacity and increase the strength of Mn-Cu alloys [11,12]. The Ce microalloying can suppress the enrichment of impurity elements (such as C and S etc.) and the formation of α-Mn particles at the grain and twin boundaries, thus improving the mobility of these boundaries and damping capacity of the alloys [2]. Therefore, an optimized Mn-Cu-AlZn-Ce system alloy was selected, and also the vacuum induction melting technology, hot rolling and annealing processing were employed to prepare the alloys with dendrites and homogeneous equiaxial grains in the present work. The microstructure and damping capacity of Mn-Cu based alloys with dendrites and equiaxial grains, especially the effects of compositional segregation on damping behavior of the alloys have been investigated and evaluated. For the convenience of description, the vacuum induction melted as well as hot rolled and annealed alloys are referred to VIM and HRA ones hereinafter, respectively. This paper will contribute to guiding the design and development of novel as-cast Mn-Cu based alloys, especially some complicated casting components used for controlling vibration and reducing noise. 2. Experimental procedures The ingot of Mn-Cu based alloy with a nominal composition of 70Mn-24.95Cu-3Al-2Zn-0.05Ce (at.%) was prepared by vacuum induction melting technology under an argon atmosphere with a purity of 99.999 vol%. Notably, the pressure of vacuum chamber had been pumped below 0.1 Pa before filling with argon, otherwise probably resulting in the formation of Mn oxides and thus degrading the damping capacity and mechanical properties of the alloys [4,13]. The alloy was melted in a magnesia crucible and then poured into a metal mold for cooling. Many rods with a size of 10 × 10 × 100 mm3 were cut at the same distance from the center axis of the cylindrical alloy ingot with a size of about Φ100 × 120 mm3 by electro-discharge machining (EDM). Some of alloy rods were put in a box-type resistance furnace and heated up at a temperature raise rate of 5 °C/min. When the temperature reached 500 °C, they were immediately taken out and hot rolled into strips on a rolling mill with a thickness reduction of 50%. Further, the hot rolled alloy strips were annealed at 850 °C for 5 h in a box-type resistance furnace and then water quenched. To obtain the spinodal decomposition microstructure and increase the fcc-fct transformation temperatures, the VIM rods and HRA strips were both aged at 430 °C for 2 and 4 h in a box-type resistance furnace and then water quenched. The selection of aging temperature and time is primarily based on the previous experimental results from our research group. In this work, the Mn-Cu based alloys aged for different time are referred to VIM2h, VIM4h, HRA2h and HRA4h ones hereinafter, respectively. The samples used for microstructural analysis were cut into 8 × 8 × 5 mm3 square blocks from the aged VIM rods and HRA strips by EDM. Then, they were ground using SiC-grit abrasive papers down to
3. Results and discussion 3.1. Microstructure of VIM and HRA alloys Figs. 1 and 2 show the BSE images and XRD patterns of the Mn-Cu based alloys upon aging treatment, respectively. It can be seen that the microstructure of the VIM alloy mainly comprises γ-MnCu dendrites with distinct compositional segregation between dendritic and interdendritic, i.e. dark and light zones with area fractions of about 73% and 27% respectively (Fig. 1a). The EDS analysis results show that the compositions of dendritic and interdendritic zones are 76.9Mn-18.6Cu3.1Al-1.4Zn and 61.3Mn-34.4Cu-2.8Al-1.5Zn (at.%) respectively, indicating the primarily Mn and Cu segregations. Furthermore, Mn is evidently predominant in the initially solidified dendritic zone compared to terminally solidified interdendritic one due to the higher melting temperature of Mn than Cu. In contrast, the microstructure of the HRA alloy is mainly composed of homogeneous equiaxial γ-MnCu grains with an approximate composition of 71.2Mn-24.3Cu-2.9Al-1.6Zn 2
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the alloys. In this work, both the (220) and (202) peaks have been fitted using quadratic polynomial function, i.e. I = M(2θ)2 + N(2θ) + L, and the corresponding coefficients and parameters calculated by XRD analysis are presented in Table 1. It can be found that the 1-c/a value of VIM alloy is higher than that of HRA one. Moreover, with increasing the aging time from 2 to 4 h, the 1-c/a value of the former alloy increases while that of the latter one deceases instead (0.0246, 0.0303, 0.0208 and 0.0066 for VIM2h, VIM4h, HRA2h and HRA4h alloys respectively). Generally, the higher 1-c/a value, i.e. larger tetragonal distortion implies more {101} twins induced by fcc-fct transformation and thus higher damping capacity of the alloy [15]. As mentioned above, the αMn phase is inclined to form in the HRA alloy especially after aging for 4 h (Figs. 2 and 3). Also considering that the HRA4h alloy has the lowest 1-c/a value, it is easily concluded that the existence of α-Mn phase certainly slows down the fcc-fct phase transformation and therefore the formation of {101} twins in this alloy to some extent [2,4].
Fig. 2. XRD patterns of Mn-Cu based alloys.
(at.%), as shown in Fig. 1b. It should be noted that the element Ce has not been determined by EDS analysis, which is mainly associated with its very small amount in the VIM and HRA alloys. The several diffraction peaks of α-Mn phase are detected only in the XRD pattern of the HRA4h alloy (Fig. 2). It is implied that the overaging which means the formation of α-Mn with increasing the aging time, prefers to occur in the plastically worked alloy instead of vacuum induction melted one. The fine α-Mn particles with sizes in nanoscale are distributed within γ-MnCu grains and also at γ-MnCu grain boundaries, as shown in Fig. 3. Furthermore, some high density dislocations generating during hot rolling are still observed around these α-Mn particles (as indicated in Fig. 3). It is therefore inferred that the distortion regions caused by high density dislocations would have provided nuclei for the formation of α-Mn phase in HRA alloy during aging. Fig. 4 shows the locally magnified XRD patterns of VIM and HRA alloys. It can be seen that the original (220) diffraction peak separates into both (220) and (202) ones in the XRD patterns of the two alloys irrespective of aging time, which means that the fcc-fct phase transformation occurs and {101} twins form [15]. The typical twinning bands with different orientations from VIM2h alloy are displayed in Fig. 5. Moreover, there is an intersection angle of about 60° between twinning bands which is very close to the intersection angle between {101} boundaries in different fct twin variants [16], calculated approximately by the following equation usually applied to fcc structure (fct and fcc structures have very small difference of lattice parameters):
cos θ =
3.2. Damping capacity of VIM and HRA alloys Fig. 6 shows the variation of the damping capacity and relative elastic modulus of VIM and HRA alloys with temperature at a strain amplitude of 5 × 10−4. It can be seen from Q−1-T curves that the VIM alloy possesses significantly higher damping capacity than HRA one irrespective of aging time. Generally, there is a phase transformation damping peak and a twin relaxation internal friction one at high and low temperature sides in the Q−1-T curves of Mn-Cu based alloys, respectively [17]. In this work, a nontypical phase transformation peak, namely a phase transformation damping platform occurs at high temperature side in the Q−1-T curves of the two VIM alloys (Fig. 6a). In contrast, there is a relatively typical phase transformation peak in the Q−1-T curve of HRA2h alloy, however which is not observed in that of HRA4h one again due to its very low phase transformation damping capacity (Fig. 6a). As mentioned in section 3.1, many fine α-Mn particles formed in HRA4h alloy slow down the fcc-fct phase transformation and the formation of {101} twins, thus leading to the lowest damping capacity of this alloy. In addition, the obviously towering twin relaxation peaks are displayed in the Q−1-T curves of VIM alloys compared to HRA ones (Fig. 6a). It should be noted that the phase transformation damping capacity of VIM2h alloy is lower than that of VIM4h one, but the twin relaxation internal frication of the former alloy is evidently higher than that of the latter one (Fig. 6a). Fig. 7 shows the variation of damping capacity of VIM and HRA alloys with strain amplitude at room temperature. It is also observed that the VIM alloys have higher damping capacity than HRA alloys over the whole strain amplitude range irrespective of aging time. Moreover, the damping capacity of both alloys decreases with increasing the aging time from 2 to 4 h to some extent. It is well known that the temperature corresponding to minimum elastic modulus in E-T curve represents the starting phase transformation one of Mn-Cu based alloys [18]. As observed in Fig. 6b, there are similar modulus valley temperatures, i.e. starting fcc-fct transformation ones in Er-T curves of VIM and HRA alloys and they raise with increase in aging time (116.1 vs 122.6 and 146.2 vs 155.0 °C for VIM2h vs HRA2h and VIM4h vs HRA4h alloys, respectively). The similar starting fcc-fct transformation temperatures are probably related to the relatively close average Mn concentrations of Mn-rich regions in the decomposition microstructures from dendritic zone for VIM alloy and equiaxial grain for HRA one, respectively. The Mn concentration of Mnrich regions can be approximately counted as 88.8 vs 89.1 and 91.0 vs 91.7 at.% for VIM2h vs HRA2h and VIM4h vs HRA4h alloys respectively by the following empirical equation [3,4]:
H1 H2 + K1 K2 + L1 L2 H12
+ K12 + L12 ⋅ H22 + K22 + L22
(1)
where θ and (HKL) are crystal plane intersection angle and crystal plane indices respectively. Notably, the separation extent of (220) diffraction peak in the XRD pattern differs from each other (Fig. 4), indicating the various axial ratios (c/a) and tetragonal distortions (1-c/a) of fct γ′-MnCu phase in
Fig. 3. Typical TEM images from HRA4h alloy, showing the fine α-Mn particles distributed within γ-MnCu grains (a) and also at γ-MnCu grain boundaries (b), and high density dislocations around these particles as well.
Ms = 1264.8cMn – 731.8
(2)
where Ms represents starting phase transformation temperature (K) and 3
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Fig. 4. Locally magnified XRD patterns of Mn-Cu based alloys, showing the separation of (220) diffraction peak into (220) and (202) ones: (a) VIM2h, (b) VIM4h, (c) HRA2h and (d) HRA4h.
cMn is Mn concentration (at.%). It should be pointed out that the fcc-fct phase transformation occurs firstly in Mn-rich region in the decomposition microstructure of dark dendritc zone and then in that of light interdendritic one, since the former zone has higher Mn concentration than the latter one in the VIM alloy (76.9 and 61.3 at.% respectively) [3,4], as shown in Fig. 1a. In other words, the two stepped fcc-fct phase transformation happens and couples to each other in the VIM alloy with obvious compositional segregation, which inevitably imposes an important influence on its damping capacity, especially phase transformation damping capacity. This illustrates the phase transformation damping platform in the Q−1T curve of the VIM alloy in comparison with the phase transformation damping peak of the HRA alloy (Fig. 6a). Furthermore, with increasing the aging time from 2 to 4 h, the starting phase transformation temperature increases due to the more sufficient spinodal decomposition [5,15], thus resulting in a higher and wider phase transformation damping platform in VIM4h alloy than in VIM2h one (Fig. 6a). On one hand, the high fcc-fct transformation temperature promotes the formation of fct {101} twins, which improves the damping capacity of Mn-Cu based alloys [19]. On the other hand, however, the axial ratio c/a of fct phase goes down, namely the tetragonal distortion (1-c/a) increases with decrease in temperature so that the {101} planes
Fig. 5. TEM image of typical twinning bands from VIM2h alloy, showing an intersection angle of about 60° between twinning bands with different orientations.
Table 1 The coefficients M, N, L in quadratic polynomial function, i.e. I = M(2θ)2 + N(2θ) + L fitted for separated (220) and (202) peaks and corresponding peak 2θ in Fig. 4, and the lattice parameters c and a, axial ratio c/a and tetragonal distortion 1-c/a of fct γ′-MnCu phase in the alloys calculated by XRD analysis. Alloy
Peak
M
N
L
Peak 2θ
c
a
c/a
1-c/a
VIM2h
(220) (202) (220) (202) (220) (202) (220) (202)
−1415.9 −493.4 −1303.8 −1305.1 −224.7 −1016.6 −7165.9 −2050.2
200537.3 70906.4 184486.5 188030.3 31966.4 146410.5 1012020.0 290620.3
−7.1E6 −2.5E6 −6.5E6 −6.8E6 −1.1E6 −5.3E6 −3.6E7 −1.0E7
70.8162 71.8549 70.7495 72.0367 71.1313 72.0099 70.6136 70.8830
3.6707
3.7634
0.9754
0.0246
3.6522
3.7665
0.9697
0.0303
3.6708
3.7489
0.9792
0.0208
3.7480
3.7728
0.9934
0.0066
VIM4h HRA2h HRA4h
4
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Fig. 6. Variation of damping capacity (a) and relative modulus (b) of Mn-Cu based alloys with temperature at a strain amplitude of 5 × 10−4.
4. Conclusions (1) The microstructure of VIM alloy mainly consists of γ-MnCu dendrites with obvious compositional segregation while that of HRA alloy primarily comprises equiaxial γ-MnCu grains. Also, some very fine α-Mn particles are formed in HRA alloy aged for 4 h. (2) The dendrtic segregation has no significant influence on the starting fcc-fct transformation temperature of Mn-Cu based alloys which is primarily related to aging time. However, the VIM alloy has larger tetragonal distortion 1-c/a than HRA one. Moreover, accompanying the raise in aging time, 1-c/a value of the former alloy increases while that of the latter one deceases. (3) The two stepped fcc-fct transformation occurs and couples to promote the formation of a phase transformation damping platform in VIM alloy compared to a phase transformation damping peak in HRA one, which is still enhanced with increase in aging time. (4) A distinctly towering twin relaxation internal frication peak is present in VIM alloy especially aged for 2 h because of the synergetic damping effect between {101} twins in the dendritic and interdendritic zones, compared to HRA one.
Fig. 7. Variation of damping capacity of Mn-Cu based alloys with strain amplitude at room temperature.
respectively in the two twinned constituent grains might become separated from each other and incoherent. In order to maintain the coherent {101} twin boundary, the initially parallel {101} planes in these two constituent grains then rotate in opposite directions. As thus, the initial orientation relationship between the two twinned constituent grains is destroyed, i.e. their relative reorientation occurs, which results in the reduction of damping capacity at lower temperatures while the twin boundary relaxation is explained by a thermal activation process [19]. In this work, owing to having higher starting fcc-fct transformation temperature, the tetragonal distortion (1-c/a) of the VIM4h alloy is larger than that of the VIM2h one (0.0303 and 0.0246 respectively, Table 1). Therefore, a bigger destruction of initial twinning relationship and a more evident relative reorientation between the two fct twinned constituent grains are present in the former alloy than in the latter one. This illustrates the obviously lower twin relaxation internal friction of VIM4h alloy than VIM2h one to some extent (Fig. 6a). As mentioned above, the starting phase transformation temperature of Mn-Cu based alloys is not significantly influenced by the compositional segregation (116.1 and 122.6 °C for VIM2h and HRA2h alloys with γ-MnCu dendrites and equiaxial γ-MnCu grains, respectively). Nevertheless, their damping capacity, especially twin relaxation internal friction is closely associated with such segregation (Figs. 6a and 7). Considering that the two stepped fcc-fct phase transformation occurs respectively in the dendritic and interdendritic zones of VIM alloy and the corresponding collaborative phase transformation damping exists (Fig. 6a), so it is easily imagined that the resulting {101} twins in these two zones would also have cooperated to improve its damping capacity. That is to say the synergetic effect of {101} twins in the two zones is probably responsible for higher twin relaxation internal friction of VIM2h alloy than HRA2h one. In addition, the relatively high density dislocations in some localized regions (the typical TEM image of some high density dislocations from HRA alloy is shown in Fig. 3) may also be another reason for degrading the damping capacity of HRA2h alloy.
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