2019 年 60 巻 11 号 p. 2298-2304
An alloy with a nominal composition of 70Mn–24.95Cu–3Al–2Zn–0.05Ce (at%) was prepared using vacuum induction melting (VIM) technology, or followed by directional solidification (DS) processing at withdrawal rates of 20 and 100 µm/s. Further, VIM, DS20 and DS100 alloys were aged at 703 K for 2 h. The microstructure, fcc-fct transformation and damping capacity of VIM and DS alloys have been investigated comparatively. The results show that the microstructure of VIM alloy mainly comprises equiaxial γ-MnCu dendrites while that of DS20 and DS100 ones is primarily composed of columnar γ-MnCu dendrites, and the directional effect of such columnar dendrite is obviously strengthened with increase in withdrawal rate. Two and three compositional segregations are present in VIM and DS alloys respectively, and fine α-Mn phase is formed in DS100 one. The starting fcc-fct transformation temperature of the alloy bears a relationship of TtVIM > TtDS20 > TtDS100. The stepped fcc-fct transformations occur and couple to promote the formation of phase transformation damping step due to compositional segregation, which is more obvious in DS alloys than in VIM one. The twin relaxation peak damping capacity of VIM and DS20 alloys is similar but evidently higher than that of DS100 one. The damping capacity of long columnar dendrite especially at 100 µm/s is also degraded due to strong grain boundary blocking effect. There exists a relationship of Q−1VIM > Q−1DS20 > Q−1DS100 for damping capacity of VIM, DS20 and DS100 alloys at room temperature over the whole strain amplitude range.
In recent decades, controlling vibration and reducing noise have been a major engineering problem that is highly concerned and urgent to be solved in many fields such as aerospace, nuclear industry, weapon equipment, automobile and rail transit etc. It is therefore very meaningful to develop high damping materials that can absorb external vibration energy and consume it via heat dissipation.1) Owing to possessing high damping capacity (specific damping capacity, SDC: 20–40%) and remarkable mechanical properties (tensile strength, σb: 490–608 MPa; elongation, δ: 20–40%), Mn–Cu based alloys have attracted much attention and possessed broad application prospects in the fields mentioned above.2,3)
It has been extensively reported that the excellent damping capacity of Mn–Cu based alloys is mainly associated with the stress induced movement of {101} twin boundaries originating from martensitic phase transformation, i.e. fcc-fct transformation.4,5) Generally, Mn–Cu based alloy has a relatively low martensitic transformation temperature which is linearly dependent on Mn content in the alloy. It has been pointed out that the fcc-fct transformation temperature of the alloy can be raised by aging at 673–873 K due to the formation of spinodal decomposition microstructure composed of Mn-rich and Cu-rich regions with sizes of 101 nm order. Correspondingly, the abundant {101} twins are formed at room temperature, thus leading to the improved damping capacity of the alloys.6–8)
Mn–Cu based alloys are generally prepared by vacuum induction melting, subsequent plastic working (hot-forging or hot-rolling) and further heat-treatment (annealing and aging).9) Recent researches have indicated that owing to the existence of compositional segregation, the direct vacuum induction melted Mn–Cu based alloys with dendrites exhibit higher damping capacity than the plastically worked and annealed ones with homogeneous grains upon proper aging treatment.10–12) This seems to provide a new idea for the design and development of Mn–Cu based alloys by the direct melting and pouring processes without subsequent plastic working and annealing. Furthermore, it is usually more appropriate for some complicated components to be produced by casting process.13)
It is well known that the fine equiaxial dendrite zone on the surface, columnar dendrite zone and central equiaxial dendrite zone are usually present in the casting ingot. Moreover, three or two dendrite zones and even only one dendrite zone can be obtained by changing or controlling casting process parameters such as alloy composition, pouring temperature, cooling conditions, physical and mechanical fields etc.14) However, the effects of different dendrite zones, exactly dendritic morphology and segregation on the formation of {101} twins induced by fcc-fct phase transformation, and damping capacity of Mn–Cu based alloys have not been clearly understood. The illustration of this problem will contribute to guiding the selection of casting process parameters such as pouring temperature and cooling rate etc. which impose a significant influence on solidification microstructure of the alloys containing dendritic morphology, grain size and segregation degree etc.
In the present work, the vacuum induction melting technology and directional solidification processing (liquid metal cooling, i.e. LMC method) were employed to prepare Mn–Cu based alloys with equiaxial and columnar dendrites, respectively. Furthermore, the growth effect of columnar dendrite was controlled by changing withdrawal rate during directional solidification. The microstructure, fcc-fct transformation and damping capacity of Mn–Cu–Al–Zn–Ce system alloys with different dendritic morphologies and segregations have been investigated and evaluated. The additions of Al, Zn and Ce are to improve the casting performance and damping capacity of the alloy according to the open literatures.2,15) For the convenience of description, the alloys in this study are referred to VIM, DS20 and DS100 ones (VIM is vacuum induction melting technology, DS20 and DS100 represent directional solidification processing at withdrawal rates of 20 and 100 µm/s respectively).
The cylindrical master alloy ingot of 6 kg in weight with a nominal composition of 70Mn–24.95Cu–3Al–2Zn–0.05Ce (at%) was prepared using vacuum induction melting technology.16) Some fine round bars used for directional solidification were cut from master alloy ingot by electro-discharge machining (EDM). After polishing and cleaning, these bars were put into a corundum crucible with a size of 12 mm in inner diameter and 140 mm in height, which was prefixed to a convex platform jointed with a stainless steel withdrawing rod. Moreover, this rod was soaked in Ga–In–Sn cooling liquid in the crystallizer of a homemade vacuum directional solidification furnace. The directional solidification experiment was conducted using LMC method at withdrawal rates of 20 and 100 µm/s respectively. The directionally solidified alloy bars had a size of about Φ12 × 100 mm3. Generally, there exist unmelted zone, initial transition zone, steady state zone and quenched solid/liquid interface zone parallel to growth direction, i.e. withdrawal direction successively in the directionally solidified bars. Notably, all the directionally solidified specimens were taken from the steady state zone of the alloy bars.
In order to obtain the spinodal decomposition microstructure and increase the fcc-fct phase transformation temperatures, the VIM, DS20 and DS100 specimens were aged at 703 K for 2 h in a box-type resistance furnace and then water quenched. The specimens used for microstructural analysis and damping measurement were cut into 7 × 7 × 10 mm3 square blocks and 1.5 × 1.5 × 50 mm3 wires by EDM, respectively. Moreover, these blocks were ground using SiC-grit abrasive papers down to 2000#, then polished and cleaned in an ultrasonic alcohol bath, and eventually blown-dry successively.16)
The constituent phases of the specimens were identified by X-ray diffraction analysis using ICDD cards (XRD, Panalytical X’Pert PRO, Cu Kα). The microstructure of the specimens 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 specimens by EDM, then mechanically ground down to 50 µm and dimpled to 10 µm, and finally ion-milled. The chemical compositions of constituent phases in the specimens were analyzed by energy dispersive X-ray spectroscopy (EDS, Inca X-sight).16)
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 specimens were measured in an internal friction instrument using inverted torsion pendulum mode. The measurement had been performed during heating from 193 to 473 K at a temperature raise rate of 2.5 K/min, and with a frequency of 1 Hz and a strain amplitude of 5 × 10−4. Notably, the Er-T curve can be used to assess fcc-fct phase transformation temperature of the alloy.9) Besides, the variation of damping capacity of the specimens with strain amplitude (Aε) from 0 to 10−3 order was obtained at room temperature with a frequency of 1 Hz.16)
Figure 1 shows the XRD patterns of VIM, DS20 and DS100 alloys. It is found that there are only diffraction peaks from γ-MnCu phase detected in the XRD patterns of VIM and DS20 alloy. In addition to γ-MnCu phase, however, α-Mn phase is present in DS100 alloy, displaying several diffraction peaks of this phase in the XRD patterns. As shown in Fig. 2, the obvious separation of (220) diffraction peak into (220) and (202) ones was detected in the locally magnified XRD patterns of VIM alloy instead of DS20 and DS100 ones, indicating larger tetragonal distortion 1-c/a of fct phase, i.e. implying more sufficient fcc-fct transformation and more {101} twins formed in the former alloy than in the latter two ones.17)
XRD patterns of Mn–Cu based alloys.
Locally magnified XRD patterns of Mn–Cu based alloys: (a) VIM, (b) DS20 and (c) DS100.
Figure 3 shows the BSE images of microstructures of VIM, DS20 and DS100 alloys and Table 1 presents the chemical compositions of their constituent phases. Combined with XRD, BSE and EDS analysis results, the microstructure of VIM alloy mainly comprises equiaxial γ-MnCu dendrites (Fig. 3(a)) while that of DS20 and DS100 ones primarily consists of columnar γ-MnCu dendrites (Figs. 3(b)–(e)). Moreover, some fine α-Mn particles with sizes in nanoscale are present in the microstructure of DS100 alloy by TEM analysis, as shown in Fig. 4. It is easily concluded that α-Mn phase is inclined to form in DS alloy especially with a higher withdrawing rate compared to VIM one. It can also be seen from the longitudinal microstructure of DS alloy that the directional growth effect of such columnar γ-MnCu dendrite is evidently enhanced with increase in withdrawal rate from 20 to 100 µm/s, as displayed in Figs. 3(c) and (e) respectively.
BSE images of VIM (a), DS20 (b, b′, c) and DS100 (d, d′, e) alloys: (b, b′, d, d′) transverse and (c, e) longitudinal microstructures. (b′) and (d′) show the details of transverse microstructures of DS20 and DS100 alloys in (b) and (d), respectively.
Typical TEM image from DS100 alloy, showing the precipitation of fine α-Mn particle.
In addition, the obvious dendritic segregation is observed in both VIM and DS alloys with two and three compositional contrasts respectively, as indicated by sites ‘1’–‘2’ for VIM, ‘3’–‘5’ for DS20 and ‘6’–‘8’ for DS100 in Fig. 3, indicating the segregations of mainly Mn and Cu instead of Al and Zn by EDS analysis (Table 1). On the whole, for VIM, DS20 and DS100 alloys, there are relatively similar Mn concentrations in their dark regions (sites ‘1’, ‘3’, ‘6’) varying in the range of 76.95–77.56 at%, gray regions (sites ‘4’, ‘7’) in the range of 65.21–65.89 at%, and also light regions (sites ‘2’, ‘5’, ‘8’) in the range of 58.77–60.58 at%, respectively. Notably, the element Ce has not been determined by EDS analysis, which is mainly attributed to its very small amount in the alloys (0.05 at%). It should be pointed out that owing to very small sizes of α-Mn particles, the compositions of dark or gray or light regions in DS100 alloy are average ones from both γ-MnCu matrix and α-Mn precipitates.
3.2 Damping capacity of Mn–Cu based alloysFigure 5 shows the variation of the relative elastic modulus and damping capacity of VIM, DS20 and DS100 alloys with temperature at a strain amplitude of 5 × 10−4. It can be seen that the temperature corresponding to valley modulus in Er-T curves, i.e. the starting fcc-fct transformation temperature of the alloy18) obviously varies and exhibits a relationship of TtVIM > TtDS20 > TtDS100 (389.1, 357.8 and 338.3 K for VIM, DS20 and DS100 alloys respectively, Fig. 5(a)). As mentioned in section 3.1, α-Mn phase tends to precipitate in DS alloy especially with a higher withdrawing rate compared to VIM one. It is easily imagined that the fine α-Mn precipitation decreases Mn concentration of γ-MnCu matrix around (very evident in DS100 alloy compared to DS20 one), which imposes a significant influence on the fcc-fct transformation temperature of the alloy depending on linearly Mn concentration.6–8) This illustrates the relationship of starting fcc-fct transformation temperature of VIM, DS20 and DS100 alloys to some extent.
Variation of relative elastic modulus (a) and damping capacity (b) of Mn–Cu based alloys with temperature at a strain amplitude of 5 × 10−4.
It should be noted that below the modulus valley temperature, the relative elastic modulus of the alloy integrally has an approximately wave-like increasing tendency with decease in temperature, which has generally not been found for Mn–Cu based alloys with homogeneous grains instead of dendrites.8) Moreover, a concave or inflection point of relative elastic modulus to temperature is observed in Er-T curve and the concave or inflection temperatures of VIM, DS20 and DS100 alloys follow a relationship of TcVIM < TcDS20 ≈ TcDS100 (as indicated by arrows in Fig. 5(a)). Considering the correlation between abnormal variation of modulus and martensitic phase transformation,18) it is concluded that the existence of concave or inflection temperature implies further fcc-fct transformation of the alloy at lower temperatures. In other words, accompanying the decrease in temperature, the stepped fcc-fct transformation has occurred in the alloys in this work, which is closely related to the compositional segregation present in their dendrites since the martensitic transformation temperature linearly depends on Mn content in the dark, gray and light regions (Fig. 3 and Table 1).6,7) It should be pointed out that this stepped fcc-fct transformation is more evident in DS alloy with three segregation regions than in VIM one with just two segregation regions. This can be confirmed by the results from Q−1-T curves in Fig. 5(b) to some extent (described below).
Generally, there is a phase transformation damping peak and a twin relaxation one at high and low temperature sides in Q−1-T curves of Mn–Cu based alloys with homogeneous grains, respectively.8) In this work, however, not a typical phase transformation damping peak but a damping step is observed in Q−1-T curves of VIM, DS20 and DS100 alloys, and the corresponding step temperature conforms to such a relationship of TsVIM > TsDS20 > TsDS100 similar to TtVIM > TtDS20 > TtDS100 (as indicated by arrows in Fig. 5(b)). Furthermore, this phase transformation damping step is steeper in the latter two alloys than in the former one. As described above, the stepped martensitic phase transformation occurs in the alloys with dendrites due to compositional segregation. It is easily imagined that the two or more stepped fcc-fct transformations certainly overlay and couple to each other, thus leading to the formation of phase transformation damping step instead of damping peak, which is still enhanced in DS alloy compared to VIM one as more obvious compositional segregation exist in the former alloy than in the latter one (three vs. two segregation regions, Fig. 3 and Table 1).
It can also be seen from Fig. 5(b) that across the phase transformation damping step, the damping capacity contributed by twin relaxation under a thermal activation,19) continues to increase with further decease in temperature and reaches a maximum value at 257.4, 250.7 and 245.9 K for VIM, DS20 and DS100 alloys respectively. And the twin relaxation peak damping capacity of VIM and DS20 alloys is relatively close but significantly higher than that of DS100 one. The typical twinning bands from VIM alloy are displayed in Fig. 6. In addition, the damping capacity of the alloys at room temperature exhibits a relationship of Q−1VIM > Q−1DS20 > Q−1DS100 similar to the variation of their starting phase transformation temperatures over the whole strain amplitude range (0–10−3 order), as shown in Fig. 7.
TEM image of typical twinning band from VIM alloy.
Variation of damping capacity of Mn–Cu based alloys with strain amplitude at room temperature.
Normally, the high phase transformation temperature promotes the fcc-fct martensitic transformation and the formation of {101} twins, resulting in the improved damping capacity of the alloy.17) Based on the relationship of TtVIM > TtDS20 > TtDS100 and XRD results about separation of (220) diffraction peak or not (Figs. 1 and 2), it is inferred that probably more {101} twins form in VIM alloy than in DS20 and especially DS100 ones, thus contributing to higher damping capacity of the former alloy than the latter two ones at room temperature.
As described in section 3.1, the equiaxial and columnar dendrites exist in VIM and DS alloys respectively, and strongly directional growth effect of columnar dendrite is obtained in DS100 alloy compared to DS20 one. And the directional solidification can eliminate transverse grain boundaries to some extent, and almost the longitudinal grain boundaries are retained instead, which is still strengthened by increasing withdrawal rate.14) It is well known that the fct {101} twins have been induced by releasing the stress caused by {110}⟨110⟩ martensitic shears (i.e. stress release mechanism). Moreover, such martensitic shear is generally blocked by grain boundaries due to the misorientation between grains and stress concentration.20) It is easily imagined that the distribution of misorientation between grains is relatively concentrated rather than scattered for the columnar dendrite with long perpendicular boundary in DS alloys compared to equiaxial dendrite with random oriented conventional boundary in VIM one.14) The concentrated distribution of misorientation between grains leads to large resistance of {110}⟨110⟩ martensitic shears and thus high stress concentration at grain boundaries due to absence of synergetic effect between grains when releasing the shear stress. In other words, the fcc-fct transformation and formation of {101} twins might be impeded or slowed down by long columnar dendrites in DS alloys compared to equiaxial dendrites in VIM one to some extent. This is partly confirmed by the fact that there is no obvious separation of (220) diffraction peak in the XRD patterns of DS alloys, meaning a low tetragonal distortion 1-c/a of fct twins. For the same reason, the relaxation movement of {101} twin boundaries is more likely to be restricted by long perpendicular boundary of columnar dendrite than by random oriented conventional one of equiaxial dendrite based on such grain boundary blocking effect. In conclusion, the dendritic morphology has a relatively important influence on fcc-fct transformation and the formation of {101} twins, and therefore the damping capacity of the alloys.
Yan et al.1) and Zhong et al.21) have indicated that the higher density of twin bands is more easily induced by fcc-fct martensitic transformation in the large grain than in the small one. It is therefore suggested that the DS20 alloy has probably more {101} twins than DS100 one due to the obviously larger average grain size of dendrites in the former alloy than in the latter one (Figs. 3(b) and (d)). Besides, the precipitation of equilibrium α-Mn phase in DS100 alloy (Figs. 1 and 4) may somewhat suppress the formation of {101} twins and hinder their movement,2) however, which has almost not taken place in VIM and DS20 alloys. All of these illustrate the damping relationship of Q−1VIM > Q−1DS20 > Q−1DS100 at room temperature and strain amplitude of 0–10−3 order.
Owing to more obvious stepped fcc-fct transformation and coupling, the synergetic effect of {101} twins formed in three segregation regions of DS20 alloy is probably greater than that in two segregation regions of VIM one. And the greater synergetic effect of {101} twins present in different segregation regions means more obviously improved damping capacity of the alloy. This illustrates that although the starting phase transformation temperature of DS20 alloy is lower than that of VIM one, implying relatively less {101} twins in the former alloy than in the latter one, they can still have similar twin relaxation peak damping capacity (Fig. 5(b)).
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).
