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Effect of alloying with Al and Cr on the microstructure, damping capacity and high-temperature oxidation behaviors of Fe–17Mn damping alloys
Journal of Alloys and Compounds xxx (xxxx) xxx
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Effect of alloying with Al and Cr on the microstructure, damping capacity and high-temperature oxidation behaviors of Fee17Mn damping alloys Guang Chen a, Benjamin Giron-Palomares b, Hongying Sun c, Hui Wang d, *, Yiyong Zhang e, Lian Duan a, Jun Wang a, Ping Cao f a
School of Mechanical Engineering, Sichuan University, Chengdu, Sichuan, 610065, China Engineering Training Center, Anyang Institute of Technology, Anyang, 455002, China School of Mechanical Engineering, Anyang Institute of Technology, Anyang, 455002, China d Science and Technology on Reactor Fuel and Materials Laboratory, Nuclear Power Institute of China, Chengdu, Sichuan, 610041, China e School of Aeronautics and Astronautic, Sichuan University, Chengdu, Sichuan, 610065, China f Zhejiang Jiuli Special Materials Technology Co., Ltd., Huzhou, Zhejiang, 313000, China b c
a r t i c l e i n f o
a b s t r a c t
Article history: Received 17 September 2019 Received in revised form 13 November 2019 Accepted 14 November 2019 Available online xxx
Fee17Mn is a commonly used damping alloy with excellent strength, wide temperature range, and low cost. Nevertheless, its poor resistance to corrosion and its low high-temperature oxidation resistance limit its application. In this study, the high-temperature oxidation resistance was enhanced by alloying with Cr and Al. The oxidation resistance was analyzed at 500 C. Besides, the effects of alloying with Cr and Al on the microstructure and damping capacity of Fee17Mn alloys were also investigated. Alloying with Cr and Al changed the Ms temperature of the alloys and affected the solid phase composition. Lower Ms temperatures produced higher g-austenite and ε-martensite phase fractions. Al had a more significant effect on the reduction of the Ms temperature than Cr, because Al sharply increased the stacking faults energy that acted as a barrier for the g / ε phase transformation. Alloying with Cr and Al decreased damping capacity at low and high amplitudes. This decrement was a result of the reduction of the stacking faults probability and the ε-martensite. At high amplitudes, the pinning of dislocations was the main factor deteriorating damping capacity. While Cr increased the weak pinning points, Al increased the strong pinning points. The oxidation kinetics obeyed an exponential function model, and alloying with Cr and Al significantly decreased the rate constant. The oxide scales of the Fee17Mn binary alloy, which easily peeled off during cooling, mainly consisted of M2O3 and MnFe2O4 with several voids. The outmost of the substrate formed a ferritic layer due to the selective oxidation of Mn at high temperatures. MnO was found at the interface between the oxide scales and the ferritic layer. Although alloying with Cr and Al was not enough to form oxide scales, Cr and Al oxides dispersed in the Mn oxide scales and enriched the top of the ferritic layer hindering the inward diffusion of oxygen. © 2019 Elsevier B.V. All rights reserved.
Keywords: Fe-17Mn alloy Alloying with Cr and Al Microstructure Damping capacity High-temperature oxidation
1. Introduction With people’s pursuit of a quiet and comfortable living environment, vibration and noise need to be controlled. It is undesirable and inadequate for conventional methods to reduce vibration in engineering design, because the size and weight must be minimized while assuring an adequate and complex vibration spectra
* Corresponding author. E-mail address: [email protected] (H. Wang).
[1]. Therefore, a kind of functional material featuring high damping capacity and good mechanical properties has been proposed [2]. Damping alloys are a technical and economic solution for vibration isolation, as well as reduction of shocks and noise. FeeMn based alloy is a new type of damping alloy designed by Korean materials researchers. When compared to nonferrous alloys like MneCu and NieTi alloys, this ferrous alloy has shown superior mechanical properties and lower cost [1e4]. Thus, it has very extensive potential applications in the fields of architecture, machinery, transportation, etc. The solid solution of the FeeMn alloy usually consists of g-austenite (fcc), ε-martensite (hcp) and a0 -
https://doi.org/10.1016/j.jallcom.2019.153035 0925-8388/© 2019 Elsevier B.V. All rights reserved.
Please cite this article as: G. Chen et al., Effect of alloying with Al and Cr on the microstructure, damping capacity and high-temperature oxidation behaviors of Fee17Mn damping alloys, Journal of Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2019.153035
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martensite (bct). Several complex phase transformations (g / ε, ε / a0 , g / a0 and ε / g) can occur among these phases [4,5]. Most scholars agree the damping mechanism of FeeMn based alloy is dominated by four factors: (ⅰ) g/ε interface boundaries, (ⅱ) ε-martensite variant boundaries, (ⅲ) stacking faults in g-austenite, and (ⅳ) stacking faults boundaries in ε-martensite. a0 -martensite is usually considered as a harmful phase for damping capacity [6]. Several researchers [5,7] have indicated that the Fee17Mn alloy exhibits its highest damping capacity in the FeeMn binary system because of the superior damping sources. In addition to varying the Mn content, many FeeMn based alloys are prepared by using two or three additional alloying elements. Many researchers studied the effects of additional alloying on the microstructure, damping capacity, corrosion resistance, and mechanical properties of the FeeMn alloys. Wen and Choi [8,9] studied the effects of carbon on the damping capacity of Fee17Mn alloys. They revealed the carbon deteriorated the damping capacity because of segregation. According to Jun et al. [10], alloying Fee23Mn with Co can increase damping capacity by increasing the length of g/ε interface (i.e. increasing the number of ε-martensite plates per unit area). Gavriljuk et al. [11] demonstrated that alloying FeeMn alloys with nitrogen provided a higher level of damping capacity by improving the mobility of interfaces. Huang et al. [12] and Jee et al. [13] pointed out that although the addition of Cr slightly reduced damping capacity, it was also a useful method to improve the corrosion resistance of FeeMn alloy. Besides, the addition of Al into FeeMn alloys can significantly decrease the g / ε transformation while improving corrosion-resistance [14]. It is known that the corrosion resistance of the FeeMn damping alloys is too low for many working conditions. The FeeMn damping alloys exhibit low corrosion resistance because of the co-existence of two or three phases, which easily form electrochemical primary cells. On the other hand, the FeeMn damping alloys can keep a high damping capacity at a wide temperature range from room temperature to 210 C [15]. Thus, it is feasible to apply FeeMn damping alloys in high temperature environment. For example, there is a strong requirement on effectively and accurately suppressing the mechanical vibration and equipment noise to ensure the quiet operation of nuclear power systems [16]. But the in-service temperature of the light water reactor is about 360e415 C, and sometimes the temperature will reach 450e500 C [17]. Therefore, it is necessary to analyze and improve the high-temperature oxidation resistance of the FeeMn damping alloys. In this paper, Fee17Mne5Cr and Fee17Mne5Cre1Al alloys were designed to comparatively study the effects of the addition of Cr and Al on the microstructure, damping capacity, and hightemperature oxidation behaviors of Fee17Mn based damping alloys. And the tests performed at 500 C was to evaluate the oxidation behaviors of materials under the critical operation temperature of nuclear reactors. The effects of alloying the FeeMn damping alloy with Cr and Al were studied by means of X-ray diffraction, SEM, DSC measurement, dynamic mechanical analyzer, and X-ray photoelectron spectroscopy. 2. Experimental methodology 2.1. Materials and solution treatment Three Fee17Mn damping alloys with different Cr and Al contents were prepared (the chemical compositions of these alloys are listed in Table 1). Ingots were prepared by arc-melting industrial pure iron and electrolytic manganese in a ZG-25A vacuum induction melting furnace. Prior to hot forging into 20-mm-thick plates at 1100 C (thermal deformation was not below 70%), the ingots were homogenized at 1150 C for 12 h. Then, the plates were
Table 1 Elemental compositions of the FeeMn damping alloys. Sample codes
Alloys compositions (wt%)
Mn
Cr
M1 M2 M3
Fee17Mn Fee17Mne5Cr Fee17Mne5Cre1Al
17.21 17.13 17.39
4.80 4.53
Al
Fe
1.03
Bal Bal Bal
exposed to a high-temperature solution treatment (at 1000 C for 1 h) with different cooling modes: water cooling (WC), air cooling (AC) and furnace cooling (FC). According to Refs. [18,19], the best solution temperature for FeeMn based damping alloys is 1000 C. The damping capacity of these three materials obtained by different cooling modes were investigated to choose the adequate cooling method. As seen in Fig. 1, damping capacity was significantly improved after solution treatment for the three cooling modes. However, the degree of improvement was as follows: water cooling was the best, air cooling provided a medium improvement, and furnace cooling was the worst. Therefore, higher cooling rates (water cooling) provide higher damping capacity and, accordingly, the alloys by water cooling were chosen for further studies. 2.2. Microstructure characterization Microstructures were observed by field emission scanning electron microscopy (FE-SEM; JSE-5900LV, Japan). Specimens were mechanically polished, electro-polished in a solution of 20% HClO4 þ 80% C2H5OH, and finally etched in a 1.2% K2S2O5 water solution. The solid phase constituents were determined by X-ray diffraction (XRD; Philips X’Pert 1, Netherlands) using Cu-Ka radiation (l ¼ 1.54 Å) in a 2q range from 35 to 85 . According to the scientific literature [20,21], the integrated intensities of ε (101), g (200) and a0 (110) were used to determine the volume fractions of ε, g and a0 in each sample. 2.3. Damping tests Damping capacity was measured by means of a forced subresonance type PL-DMTA Mk II dynamic mechanical analyzer with a size of 50 mm 2 mm 1 mm. The tests were carried out at 30 C with a 1 Hz vibration frequency and strain amplitude ranging from 0 to 1.4 103. A second damping measurement was also performed on the samples to examine their damping stability. Thus, the results of the damping capacity represented the average value of two tests. Besides, before the testing, we have checked the stability by repeating tests three times. 2.4. High-temperature oxidation tests Prior to oxidation, the surfaces of the samples were mechanically polished by using granulometric sandpapers (320, 600, and 1000 mesh). Then, each specimen was cleaned with acetone to ensure a clean and flat surface, and the dimensions of the samples were recorded by a spiral micrometer. Finally, the hightemperature oxidation tests were conducted at 500 C for 35 h in air atmosphere with a furnace. Every 5 h, three samples of each alloy were weighed, and the average weight gained by oxidation was determined intervals of 5 h. Weight gain rate can be described by the weight change per square centimeter of the specimen per minute. Surface and sectional micrographs of the high-temperature oxidized samples were observed by FE-SEM (JSE-5900LV, Japan) and energy-dispersive X-ray spectrometry (EDS). The structure of the oxide scales was analyzed by grazing incidence XRD (XRD;
Please cite this article as: G. Chen et al., Effect of alloying with Al and Cr on the microstructure, damping capacity and high-temperature oxidation behaviors of Fee17Mn damping alloys, Journal of Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2019.153035
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Fig. 1. Effects of the solution treatment cooling modes on the damping capacity for the damping alloys: (a) M1, (b) M2, (c) M3.
Philips X’Pert 1, Netherlands) with an angle of 2 , 0.04 scan step size, using Cu-Ka radiation (l ¼ 1.54 Å), and 2q range from 25 to 95 . X-ray photoelectron spectroscopy (XPS; AXIS Ultra DLD, Kratos) was also used to investigate the structure of the oxide scales. 3. Results and discussion
exhibit several lenticular a0 -martensite that is mostly embedded in ε-martensite because of the preferential orientation [22] for the formation of a0 phase. In accordance to the Burgers and the Kurdjumov-Sachs (KeS) orientation relationships [23], small a0 islands are oriented with respect to the ε variants and the g matrix. These orientation relationships are characterized by the parallelism of a close-packed plane and a close-packed direction (i.e.
3.1. Microstructure analysis As shown in Fig. 2, the XRD results indicate that several variants of martensite formed in g-austenite grains after heat treatment. During the solution treatment at 1000 C, high density of vacancies formed in the g-austenite. When the alloys were directly water cooled to room temperature, most of the vacancies were reversed by the high cooling rate. The vacancies provided many favorable sites for the nucleation of martensite during cooling. The intensity of the peaks in the XRD patterns were clearly affected by the contents of Cr and Al, which induced different solid phases. For example, g (111) and a’ (200) diffraction peaks appeared in the pattern for M3 (Fee17Mne5Cre1Al) alloys at 43.563 and 64.442 , respectively. The contents of g and a’ phases in M3 were higher than those in M1 and M2 alloys. Fig. 3 shows the SEM morphologies for the damping alloys after solution treatment. General and detail views are shown for each kind of alloy. As shown in Figs. 3(a), (b), (d), and (e), the microstructure of M1 and M2 alloys shows only ε-martensite and gaustenite plates, because the content of a0 phase is too small to be observed. The M1 alloy seems to have several thin and short ε-martensite embedded in g-austenite. Apparently, Fig. 3(c) and (f)
Fig. 2. XRD patterns of the M1, M2 and M3 damping alloys.
Please cite this article as: G. Chen et al., Effect of alloying with Al and Cr on the microstructure, damping capacity and high-temperature oxidation behaviors of Fee17Mn damping alloys, Journal of Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2019.153035
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Fig. 3. SEM morphologies of the Fee17Mn based damping alloys: (a, d) M1, (b, e) M2, (c, f) M3; (aec) general view, (def) detail view.
ð111Þg k ð110Þa’ and ½1120ε k ½111a’). In these alloys, many ε-martensite variants are parallel to each other, and this phenomenon can be explained by the next famous Shoji-Nishiyama (SeN) orientation relationships [24]: ð111Þg k ð0001Þε and ð110Þg k ð2110Þε. Therefore, Al has bigger effects on the microstructure than Cr. To further study the effects of Cr and Al on the microstructure, the volume fractions of g, ε, and a0 phases in the three alloys were calculated based on XRD patterns, and the results are listed in Table 2. These results demonstrate that alloying with only Cr (M2 alloy) has few effects on the microstructure: the volume fraction of ε-martensite decreased by 3.5%, while g-austenite and a0 phase increased by 2.1% and 1.4%, respectively. Furthermore, the volume fraction of ε-martensite in the M3 alloy decreased to 70.9%, while the g-austenite and a0 phase increased by 17.3% and 11.8%, respectively. Therefore, alloying with only Cr decreased the ε-martensite and increased g and a’ phase. Alloying with both, Al and Cr, significantly promoted the change of the phase contents. The effects of alloying elements on the formation of thermalinduced ε-martensite during solution treatment had been studied by Kim at al [25]. Their results showed that the volume fraction of thermal-induced ε-martensite notably decreased by introducing carbon and a consequent reduction of the Ms temperature. Therefore, to confirm a similar effect of Cr and Al in the FeeMn based alloys, the transformation temperatures (Ms, As, and Af) of the M1 to M3 alloys were determined by DSC measurements. As shown in Fig. 4, alloying with only Cr provoked minor decrements on the transformation temperatures. This agrees with the experimental results of Guerrero et al. [26]. Surprisingly, the transformation temperatures of M3 decreased sharply because of the
Table 2 Volume fractions of ε-martensite, g-austenite and a0 -martensite in the M1, M2 and M3 damping alloys. Samples
e-martensite
g-austenite
a0 -martensite
M1 M2 M3
89.1 85.6 70.9
9.4 11.5 17.3
1.5 2.9 11.8
joint addition of Al and Cr. Besides, the decrement of Ms was much larger than that of As, thereby expanding the temperature difference between As and Ms. Kim et al. [25] pointed out the difference between As and Ms is related to the driving force for the transformation. A larger temperature difference between As and Ms means that higher energy needed to overcome the transformation barriers. The driving force is strongly related to the martensitic transformation temperatures. And the martensitic transformation temperatures depend on the compositions of alloys. Thus, the driving force also depends on the compositions of alloys indirectly. Apparently, the driving force has a stronger dependence on Al content than on Cr content. The transformation barriers in FeeMn based alloys mostly result from the stacking fault energy (SFE), because the stacking faults are the nucleation sites for ε-martensite. The SFE of Fee17Mn alloys increased gradually with the increment of Cr content (0-12 wt%) [27,28] but increased sharply with the increment of Al content [29e31]. Therefore, the SFEs in M1, M2, and M3 keep the next relationship: SFEM1
Please cite this article as: G. Chen et al., Effect of alloying with Al and Cr on the microstructure, damping capacity and high-temperature oxidation behaviors of Fee17Mn damping alloys, Journal of Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2019.153035
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Fig. 4. MS, AS and AF temperatures measured by DSC method for: (a) M1 alloy, (b) M2 alloy, and (c) M3 alloy.
alloys increased with the increase of amplitude, but the increment was less significant for high amplitudes, especially in the M3 alloy. Alloying with Cr or Al had negative effects on the damping capacity of the FeeMn binary alloys. Alloying with both, Al and Cr, further decreased the damping capacity. For comparison, the damping capacities at strain amplitudes of 3 104 and 1.1 103 are shown in Fig. 5(b). At the low strain amplitude of 3 104, the damping capacity decreased by 21% after alloying with Cr and decreased by 44% after alloying with both Al and Cr. The damping capacity decreased by 8% after alloying with Cr and decreased by about 24% after alloying with both Al and Cr at the high strain amplitude of 1.1 103. Therefore, the effects of alloying with Cr and Al on damping capacity are different at low and high strain amplitudes. To further study the effects of alloying with Cr and Al on the damping capacity of FeeMn alloys, the classic Granatoe-Lücke (G-L) dislocation pinning model was used. According to the G-L model [35], the total damping (Q1), consisting of the strain-amplitudeindependent part Q1 0 and the strain amplitude-dependent part Q1 h , can be expressed as follows:
Q
1
¼ Q 1 0
Q 1 0 ¼
þ
Q 1 h
BLL4C u 36Gb2
Q 1 h ¼ ðC1 = εÞexpð C2 = εÞ
(1)
(2)
(3)
C1 ¼
rFB L3N 6bEL2C
C2 ¼
FB bELC
(4)
where B is the damping constant for dislocation, L is the total length of dislocation line per unit volume, u is the angular frequency, G is the shear modulus, b is the Burger’s vector of dislocations, FB is the binding force between dislocations and weak pinning points, r is the total length of dislocation line per unit volume that participated in the breakaway process, E is the Young modulus, LC is the average dislocation distance between weak pinning points, and LN is the average dislocation distance between the strong pinning points. The unpinning of dislocations is the main factor affecting damping capacity at high amplitudes. By logarithmic operations, Equation (3) can be expressed by the following equation:
C2 þ lnC1 ln Q 1 h *ε ¼ ε
(5)
where C2 and lnC1 are the slope and intercept in the G-L plots, respectively. To obtain a high Q1 value, the value of C1 should be large, while the value of C2 should be small. Fig. 5(c) illustrates the relationship between 1/ε and ln(Q1 h *ε) in the damping alloys. The linear fitting relationship between 1 ln(Q1 is well satisfied for values above the critical value h *ε) and ε (εCr ¼ 4 104), indicating that the damping capacity of the FeeMn samples can be interpreted by the G-L model when strain
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Fig. 5. Damping capacities for M1 to M3 alloys: (a) damping capacities at different strain amplitudes, (b) comparison of damping capacities at strain amplitudes of 3 104 and 1.1 103, (c) ln(Q1*ε) vs 1/ε with a linear fitted results inset, and (d) Stocking faults probability (S.F.P.) in g-austenite and ε-martensite.
amplitude only beyond εcr. The linear fitting coefficients (C1 and C2) are shown in the inset of Fig. 5(c). After alloying with both Al and Cr, the value of C2 increased with a decrease of Lc, indicating alloying with both Al and Cr shortened the average dislocation distance between weak pinning points. On the other hand, LN is the main parameter determining C1. Alloying with both Cr and Al increased the strong pinning points, while the average dislocation distance between the strong pinning points decreased. Alloying with Cr had a big effect on C2, and alloying with Al had a big effect on C1. The reason is that alloying with Cr mainly increases the weak pinning points, but alloying with Al mainly increases the strong pinning points. When the strain amplitude exceeds a critical value, Shockley partial dislocations break away from weak pinning points, rapidly increasing the damping capacity of the alloy. Thus, alloying with Cr has smaller effects on the decrease of damping capacity for high amplitudes above 4 104. Conversely, the damping at low strain amplitudes mainly results from the amount of stacking faults and phase fraction in alloys. The stocking faults probability (S.F.P.) in ε-martensite and g-austenite for M1 to M3 alloys were calculated (see Fig. 5(d)) based on Warren’s theory and peak-shift method. A detailed description of the calculation procedures can be obtained in Refs. [12,36]. The S.F.P. in ε-martensite and g-austenite slightly decreased after alloying with Cr, but significantly decreased after alloying with both Cr and Al. The reason is that SFE increased slightly in alloying with 5 wt% Cr, while sharply increased with the addition of Al. S.P.F. is inversely proportional to SFE. Besides, many researchers agree that the damping capacity increases with the increment of the amount of martensite [37,38]; accordingly, the fraction of ε-martensite decreased after alloying with both Cr and Al (check the phase
fraction in Table 2). This led to the area reduction of g/ε boundaries. 3.3. High-temperature oxidation resistance 3.3.1. Oxidation kinetics The weight gain versus time plots for the three different alloy compositions oxidized in air at 500 C for 0e35 h are shown in Fig. 6(a). Weight gain always increased with the increase of oxidation time. At the early time of the oxidation, weight gain increased rapidly. Then, the increment of weight gain became very small after a transition period. For example, the weight gain plots of M2 and M3 alloys became almost horizontal at the end of the oxidation procedure. The weight gain rates (WGR) at 5 h, 20 h, and 35 h are shown in Fig. 6(b). The weight gain rates of the damping alloys keep the next relationship for a particular oxidation time: WGRM1 > WGRM2 > WGRM3. Furthermore, weight gain rate decreased with increasing oxidation time for all the alloys. Thus, the kinetics plots can be divided into 3 stages: rapid oxidation at the initial stage (Stage A), oxidation at the intermediate stage (Stage B), and slow oxidation at the last stage (Stage C). It is well known that the process of oxidation is the diffusion of metal ions in and out from the bulk in response to the O2 chemical potential, forming sometimes a complex intermediate phase on the surface [39]. Stage A is the period in which the oxides grew rapidly on the samples surface. As time went on, the oxides formed oxide scales that decreased the speed of new oxides formation (i.e. Stage B). During Stage C, the oxide scales became thick and compact, preventing oxygen from diffusing inward. The M1 alloy only showed Stage A (0 he15 h) and the Stage B (15 he35 h), which means a poor high-temperature oxidation
Please cite this article as: G. Chen et al., Effect of alloying with Al and Cr on the microstructure, damping capacity and high-temperature oxidation behaviors of Fee17Mn damping alloys, Journal of Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2019.153035
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Fig. 6. Alloys plots for: (a) high-temperature oxidation kinetics at 500 C, and (b) weight gain rates for 5 h, 20 h, and 35 h.
resistance. The Stage A and B for the M2 alloy occurred at the intervals 0 he10 h and 10 he25 h, respectively. These two stages are shorter than those for the M1 alloy. In addition, Stage C appears after oxidizing in the interval from 25 h to 35 h. This might be due to Cr enhancing the potential of the substrate and the protective Cr oxides formed on the surface. The oxidation resistance of the FeeMn alloy was further improved after the addition of Al. The M3 alloy showed shorter Stage A (0 he5 h) and longer Stage B (5 he20 h) and Stage C (20 he35 h) than those of the M1 and M2 alloys. This indicates that the joint addition of Cr and Al is beneficial for the high-temperature oxidation resistance. The relationship between weight gain and time can be fitted by the next exponential function model:
y ¼ kp ½1 expðt = bÞ
(6)
where kp is the oxidation rate constant and b is the time constant. The fitting results are shown in the curves of Fig. 6(a). The rate constants for M1, M2, and M3 alloys are 3.37 104 g cm2 h1, 1.78 104 g cm2 h1, and 1.36 104 g cm2 h1, respectively. These indicate that the high-temperature oxidation resistance was improved by the sole addition of Cr, and further increased by the simultaneous addition of both Cr and Al. 3.3.2. Oxidation scales morphology Fig. 7 shows the microstructure and composition of the damping alloys surface after oxidation at 500 C for 20 h. As shown in Fig. 8(a), the oxide scales in M1 were severely broken and peeled off, which spontaneously occurred during the air cooling from 500 C to room temperature. The detail view of the edge of the broken oxide scales is shown in Fig. 7(b). There are many cracks in the oxide scales and the edges of the broken oxide scales are neat and sharp. This can be explained as follows: the thermal expansion and contraction coefficients of the oxide film and the base layer are different. Thus high thermal stresses generated, breaking the oxide scales [40]. The element contents of the oxide scales (area M1-1) and the peeled regions (area M1-2) are compared by EDS in Fig. 7(c). The oxide scales of the M1 alloys are mainly iron and manganese oxides. The oxygen content of the peeled regions is about 2.5 times lower than that of the oxide scales, showing that the oxygen content decreases from the outside to the inside. Fig. 7(d) is the general view of the M2 alloy samples surface showing a few light color peeled regions. A magnification of the red square in Fig. 7(d) shows a clear view of the oxide scales and peeling (see Fig. 7(e)). The oxide scales revealed an uneven and warped structure. The peeling exhibited high iron content and about 20 at% oxygen, indicating the peeling was a result of the thermal stress. After alloying with Cr, a reduction of the peeling of oxide scales was observed in the surface morphologies.
Furthermore, the peeling exhibited about 5.6 at% Cr, but the surface of the oxide scales had only 0.86 at% Cr. The microstructure of the oxide scales was covered with an oxide scale of small nodules ranging in sizes from 1 mm to 2 mm. Yuan et al. [41] observed a similar microstructure of the oxide scales in Fee25Mne3Cre3Al0.3C-0.01 N austenitic steel during a high temperature oxidation at 700 C for 8 h. This microstructure was attributed to the Cr inside the FeeMn oxides [41]. The microstructure of the M3 alloy sample surface after oxidation is shown in Fig. 8(g) and (h). No peeling occurred, and the surface of the oxide scales were composed of numbers of oxide clusters. EDS results revealed that the alloy oxide was mainly composed of iron and manganese oxides with small amounts of chromium and aluminum oxides. During the oxidation process, the oxidation of the active metal elements (Mn, Cr, and Al) preceded the oxidation of the inert metal element (Fe). According to Wang et al. [42], the M3 alloy studied in the present work had insufficient aluminum content to form continuous alumina scales. After the depletion of aluminum, diffusion of manganese towards the surface layer took place, and manganese oxide was predominantly formed leading to oxide nodules. To confirm the phase composition of the surface oxide, GIXRD analysis was performed on the surface of the damping alloys. The GIXRD results for the alloys oxidized in an air atmosphere at 500 C for 20 h are shown in Fig. 8(a). The phases were identified by matching the patterns with powder crystallographic data (PDF cards: 87e0721 (a-Ferrite), 77e2363 (MnO), 73e1964 (MnFe2O4), and 78e0390 (Mn2O3)). For the M1 alloy, the diffraction peaks corresponded with aFerrite and the oxides: Mn2O3, MnO (halite), and MnFe2O4 (spinel). The pattern of the M1 alloy exhibited the highest intensity for the a-Ferrite (110) peak, illustrating that the surface mainly contains aFerrite. This was caused by the peeling off oxides and the exposure of the matrix. Ferritic layer was likely produced on the surface of the FeeMn damping alloys at high temperatures, because the vapor pressure of Mn is four orders of magnitude higher than that of Fe. The patterns for the M2 and M3 alloys indicated the samples surface mainly contained Mn2O3 and MnFe2O4. A low intensity of aFerrite (110) was also identified. This low intensity emission probably proceeded from too thin oxide scales or small oxide peeled regions. On the other hand, the M2 and M3 alloys exhibited the highest intensity peak (222) for Mn2O3, indicating that the samples surface mainly contained Mn2O3. Because the formation of the Mn2O3 oxide requires very high oxygen partial pressures [43], the Mn2O3 phase was mainly distributed in the outmost part of the oxide scales. XPS was also applied to analyze the chemical composition and state of Mn on the surface of the samples after oxidation (see Fig. 8(b)). The Mn2p3/2 binding energy values for Mn2þ, Mn3þ, and
Please cite this article as: G. Chen et al., Effect of alloying with Al and Cr on the microstructure, damping capacity and high-temperature oxidation behaviors of Fee17Mn damping alloys, Journal of Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2019.153035
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Fig. 7. Surface morphologies and composition of the damping alloys surface after oxidation at 500 C for 20 h: (a, b) M1 surface morphologies, (c) EDS results for areas M1-1 and M1-2, (d, e) M2 surface morphologies, (f) EDS results for areas M2-1 and M2-2, (g, h) M3 surface morphologies, and (i) EDS results for Area M3.
Fig. 8. Results for the damping alloys after oxidation at 500 C for 20 h: (a) Grazing incidence XRD (GIXRD) patterns, and (b) Mn 2p3/2 XPS spectra.
Please cite this article as: G. Chen et al., Effect of alloying with Al and Cr on the microstructure, damping capacity and high-temperature oxidation behaviors of Fee17Mn damping alloys, Journal of Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2019.153035
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Mn4þ [44] were located at the intervals of 640.7 eV/-640.8 eV, 641.7 eV/-642.3 eV and 643.1 eV/-643.5 eV, respectively. These three valence states of manganese indicate three oxide compounds: Mn2þ reflects MnO, Mn3þ represents Mn2O3, and Mn4þ indicates MnFe2O4. Obviously, the M1 alloy sample surface mainly contained Mn2þ ions, and the amount of Mn2þ ions in M1 was more than that in the M2 or M3 alloys. This indicates the oxide of the M1 sample surface is MnO, which agrees with the results obtained by GIXRD. Fig. 9 shows the cross-sectional SEM back-scattering images and EDS line scan profiles of the damping alloys after oxidation at 500 C in an air atmosphere for 20 h. As shown in Fig. 9(a), no oxide scales were found adhered to the M1 alloy substrate, but some oxide regions appeared in the top region of the substrate. This proves that the oxygen penetrated the voids containing Mn oxide scales [45] and diffused into the substrate through the phase boundaries. Different phases possess different potentials, and therefore phase interfaces are prone to oxidation-reduction reactions [46]. As shown in Fig. 9(d), the oxide regions are rich in Mn and O. The yellow line scan results illustrate the ferritic layer is Mndepleted and about 2 mm wide. As shown in Fig. 9(b), some oxide scales are adhered to the M2 alloy substrate. The oxide scales were rich in Mn and O with contents of Cr and Fe gradually decreasing from the substrate to the surface (observe Fig. 10(e)). The ferritic layer width decreased to about 1.8 mm after alloying with Cr, and no internal oxidation occurred. The Cr might have enhanced the potential of the substrate and improved the potential difference of the phases [47]; therefore, the diffusion of oxygen into the substrate was difficult. The cross-sectional microstructure of the M3 alloy sample after oxidation is shown in Fig. 9(c). Homogenous and compact thin oxide scales were found to be well adhered to the substrate, indicating that their adhesive capacity was improved by adding Al. The ferritic layer of the M3 alloy sample was about 1 mm wide (see Fig. 9(f)). Although Cr and Al generally decrease from the substrate to the oxide scales, a Cr and Al rich region can be observed in the oxide scale near the substrate ferritic layer (check the orange dotted circle in Fig. 9(f)). Accordingly, it can be inferred that Cr and Al
9
Fig. 10. Schematic illustrations of the oxide scales structure: (a) M1, (b) M2, and (c) M3.
oxides assembled in the interface between the oxide scales and the alloy substrate. These interfacial oxides in-between oxide scales and a-ferritic layer are in agreement with the results of Duh et al. [48]. The Cr and Al oxides used the oxygen diffused from the Mn oxide scale and formed a compact layer making the oxygen diffusion into the substrate more difficult. Moreover, less oxygen to oxidize the Mn in the substrate made the Mn diffusion slower and had a negative impact on the growth of the oxide scales. 3.3.3. Oxidation mechanism To illustrate the effects of alloying with Cr and Al on the hightemperature oxidation resistance of the Fee17Mn based alloys, schematic illustrations of the oxide scales of the damping alloys are depicted in Fig. 10. As shown in Fig. 10(a), the outmost oxide scales of the alloy M1 mainly consist of Mn2O3 and MnFe2O4 with several
Fig. 9. Cross-sectional SEM back-scattering images and EDS line scan profiles for the damping alloys after oxidation at 500 C in an air atmosphere for 20 h: (aec) Cross sectional morphologies; and (def) element distribution for the locations shown by yellow lines a, b, and c. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)
Please cite this article as: G. Chen et al., Effect of alloying with Al and Cr on the microstructure, damping capacity and high-temperature oxidation behaviors of Fee17Mn damping alloys, Journal of Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2019.153035
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G. Chen et al. / Journal of Alloys and Compounds xxx (xxxx) xxx
scattered voids. This was caused by the much higher growth rate of Mn2O3. Nevertheless, the volume ratio between the Mn2O3 scale and the substrate is 0.48, far lower than 1 [45]. The EDS (Fig. 7(c)) and XRD (Fig. 8(a)) results of the outmost oxide layer also showed MnFe2O4. The oxide scales formed by Mn2O3 and MnFe2O4 were easy to peel off during cooling due to the different thermal expansion and contraction coefficients of the oxide scales and the substrate causing high thermal stresses. Although Mn2O3 is a ptype semiconductor [49], the rapid growth and voids in the oxide scales provoked a long initial rapid stage. Accordingly, the amount of the voids in the cracks decreased with the increase of oxidation time. The weight gain rate decreased and the medium stable oxidation stage started (see Stage B in Fig. 6(a)). Under these circumstances, MnO formed at the interface between the oxide scales and the substrate due to the low oxygen partial pressure. Besides, selective oxidation of Mn also took place at the top of the substrate. As shown in Fig. 10(b), the peeling phenomenon of oxide scales was improved by alloying with Cr. The Cr oxides were present all over the oxide scales with a very high content at the interface between the oxide scales and the Mn depleted region (see EDS results in Figs. 7(f) and 9(e)). Cr oxides decreased the voids and cracks in the oxide scales, making more difficult for the oxygen to diffuse inward. On the other hand, the high content of Cr oxides at the interface also hindered the diffusion of oxygen and Mn. Therefore, this caused the earlier start of the slow oxidation stage (see Stage C in Fig. 6(a)) and the improvement of the oxidation resistance. As shown in Fig. 10(c), the oxide scales hardened and peeling was eradicated after alloying with Cr and Al. This is beneficial for the high-temperature oxidation resistance. The Al oxides behaved as the Cr oxides (i.e. they were present all over the oxide scales with a very high content in the interface between the oxide scales and the Mn depleted region). This hardened the inward oxygen diffusion in a short period of time. Because the diffusion process was hindered by oxides, the oxidation process in the Fee17Mn damping alloys followed an exponential function. 4. Conclusions The effects of alloying with Cr and Al on the high-temperature oxidation resistance of Fee17Mn damping alloys were investigated at 500 C. In addition, the effects of the alloying element on the microstructure and damping capacity of such damping alloys were also studied. Based on the results and discussion, it can be concluded that: (1) Alloying with both Cr and Al changed the Ms temperature of the alloys and affected the solid phase compositions. Lower Ms temperatures produced higher g-austenite and ε-martensite phase fractions. Al had a more significant effect on the reduction of the Ms temperature than Cr, because Al sharply increased the stacking faults energy that acted as a barrier for the g / ε phase transformation. (2) Alloying with both Cr and Al decreased damping capacity at low and high amplitudes. This decrement was a result of the reduction of the stacking faults probability and the ε-martensite. At high amplitudes, the pinning of dislocations was the main factor deteriorating damping capacity. While Cr increased the weak pinning points, Al increased the strong pinning points. Therefore, alloying Al decreased the damping capacity more heavily at high amplitudes. (3) Alloying with both Cr and Al significantly improved the hightemperature oxidation resistance of Fee17Mn damping alloys. This alloying procedure also reduced the thermal stresses in the oxide, and consequently the peeling during cooling was reduced. Alloying with both Cr and Al can hinder
the inward oxygen diffusion by reducing voids and forming an oxide layer at the interface between the oxide scales and the substrate.
Author contribution statement Guang Chen: Conceptualization, Software, Investigation, Writing - Original Draft, Visualization. Benjamin Giron-Palomares: Writing - Review & Editing. Hongying Sun: Formal analysis, Resources. Hui Wang: Conceptualization, Methodology, Resources, Supervision, Project administration, Funding acquisition. Yiyong Zhang: Software. Lian Duan: Resources, Data Curation. Jun Wang: Resources. Ping Cao: Methodology, Resources. Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgements This work is supported by the National Natural Science Foundation of China (51971207 and 51801194). Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi.org/10.1016/j.jallcom.2019.153035. References [1] S.-H. Baik, J.-C. Kim, D.-W. Han, T.-H. Kim, J.-H. Back, Y.-K. Lee, FeeMn martensitic alloys for control of noise and vibration in engineering applications, Mater. Sci. Eng. A 438e440 (2006) 1101e1105. [2] D.D.L. Chung, Review: materials for vibration damping, J. Mater. Sci. 36 (2001) 5733e5737. [3] S.-H. Baik, High damping FeeMn martensitic alloys for engineering applications, Nucl. Eng. Des. 198 (2000) 241e252. [4] K.K. Jee, W.Y. Jang, S.H. Baik, M.C. Shin, C.S. Choi, Damping capacity in Fe-Mn based alloys, Scr. Mater. 37 (1997) 943e948. [5] J.B. Seol, J.G. Kim, S.H. Na, C.G. Park, H.S. Kim, Deformation rate controls atomic-scale dynamic strain aging and phase transformation in high Mn TRIP steels, Acta Mater. 131 (2017) 187e196. [6] H. Wang, H. Wang, R. Zhang, R. Liu, Y. Xu, R. Tang, Effect of high strain amplitude and pre-deformation on damping property of Fe-Mn alloy, J. Alloy. Comp. 770 (2019) 252e256. [7] Y.K. Lee, J.H. Jun, C.S. Choi, Damping capacity in Fe-Mn binary alloys, ISIJ Int. 37 (1997) 1023e1030. [8] W.S. Choi, B.C. De Cooman, Effect of carbon on the damping capacity and mechanical properties of thermally trained Fe-Mn based high damping alloys, Mater. Sci. Eng. A 700 (2017) 641e648. [9] Y.H. Wen, H.X. Xiao, H.B. Peng, N. Li, D. Raabe, Relationship between damping capacity and variations of vacancies concentration and segregation of carbon atom in an Fe-Mn alloy, Metall. Mater. Trans. A Phys. Metall. Mater. Sci. 46A (2015) 4828e4833. [10] J.H. Jun, D.K. Kong, C.S. Choi, The influence of Co on damping capacity of FeeMneCo alloys, Mater. Res. Bull. 33 (1998) 1419e1425. [11] V.G. Gavriljuk, P.G. Yakovenko, K. Ullakko, Influence of nitrogen on vibration damping and mechanical properties of Fe-Mn alloys, Scr. Mater. 38 (1998) 931e935. [12] S.K. Huang, N. Li, Y.H. Wen, J. Teng, S. Ding, Y.G. Xu, Effect of Si and Cr on stacking fault probability and damping capacity of FeeMn alloy, Mater. Sci. Eng. A 479 (2008) 223e228. [13] K.K. Jee, W.Y. Jang, S.H. Baik, M.C. Shin, Damping mechanism and application of Fe-Mn based alloys, Mater. Sci. Eng. A 273e275 (1999) 538e542. [14] L. Xing, Q. Zuoxiang, Z. Yansheng, W. Xingyu, L. Fengxian, D. Bingzhe, H. Zhuangqi, Study of the paramagnetic-antiferromagnetic transition and the g / ε martensitic transformation in Fe-Mn alloys, J. Mater. Sci. 35 (2000) 5597e5603.
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Please cite this article as: G. Chen et al., Effect of alloying with Al and Cr on the microstructure, damping capacity and high-temperature oxidation behaviors of Fee17Mn damping alloys, Journal of Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2019.153035
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Effect of alloying with Al and Cr on the microstructure, damping capacity and high-temperature oxidation behaviors of Fee17Mn damping alloys Guang Chen a, Benjamin Giron-Palomares b, Hongying Sun c, Hui Wang d, *, Yiyong Zhang e, Lian Duan a, Jun Wang a, Ping Cao f a
School of Mechanical Engineering, Sichuan University, Chengdu, Sichuan, 610065, China Engineering Training Center, Anyang Institute of Technology, Anyang, 455002, China School of Mechanical Engineering, Anyang Institute of Technology, Anyang, 455002, China d Science and Technology on Reactor Fuel and Materials Laboratory, Nuclear Power Institute of China, Chengdu, Sichuan, 610041, China e School of Aeronautics and Astronautic, Sichuan University, Chengdu, Sichuan, 610065, China f Zhejiang Jiuli Special Materials Technology Co., Ltd., Huzhou, Zhejiang, 313000, China b c
a r t i c l e i n f o
a b s t r a c t
Article history: Received 17 September 2019 Received in revised form 13 November 2019 Accepted 14 November 2019 Available online xxx
Fee17Mn is a commonly used damping alloy with excellent strength, wide temperature range, and low cost. Nevertheless, its poor resistance to corrosion and its low high-temperature oxidation resistance limit its application. In this study, the high-temperature oxidation resistance was enhanced by alloying with Cr and Al. The oxidation resistance was analyzed at 500 C. Besides, the effects of alloying with Cr and Al on the microstructure and damping capacity of Fee17Mn alloys were also investigated. Alloying with Cr and Al changed the Ms temperature of the alloys and affected the solid phase composition. Lower Ms temperatures produced higher g-austenite and ε-martensite phase fractions. Al had a more significant effect on the reduction of the Ms temperature than Cr, because Al sharply increased the stacking faults energy that acted as a barrier for the g / ε phase transformation. Alloying with Cr and Al decreased damping capacity at low and high amplitudes. This decrement was a result of the reduction of the stacking faults probability and the ε-martensite. At high amplitudes, the pinning of dislocations was the main factor deteriorating damping capacity. While Cr increased the weak pinning points, Al increased the strong pinning points. The oxidation kinetics obeyed an exponential function model, and alloying with Cr and Al significantly decreased the rate constant. The oxide scales of the Fee17Mn binary alloy, which easily peeled off during cooling, mainly consisted of M2O3 and MnFe2O4 with several voids. The outmost of the substrate formed a ferritic layer due to the selective oxidation of Mn at high temperatures. MnO was found at the interface between the oxide scales and the ferritic layer. Although alloying with Cr and Al was not enough to form oxide scales, Cr and Al oxides dispersed in the Mn oxide scales and enriched the top of the ferritic layer hindering the inward diffusion of oxygen. © 2019 Elsevier B.V. All rights reserved.
Keywords: Fe-17Mn alloy Alloying with Cr and Al Microstructure Damping capacity High-temperature oxidation
1. Introduction With people’s pursuit of a quiet and comfortable living environment, vibration and noise need to be controlled. It is undesirable and inadequate for conventional methods to reduce vibration in engineering design, because the size and weight must be minimized while assuring an adequate and complex vibration spectra
* Corresponding author. E-mail address: [email protected] (H. Wang).
[1]. Therefore, a kind of functional material featuring high damping capacity and good mechanical properties has been proposed [2]. Damping alloys are a technical and economic solution for vibration isolation, as well as reduction of shocks and noise. FeeMn based alloy is a new type of damping alloy designed by Korean materials researchers. When compared to nonferrous alloys like MneCu and NieTi alloys, this ferrous alloy has shown superior mechanical properties and lower cost [1e4]. Thus, it has very extensive potential applications in the fields of architecture, machinery, transportation, etc. The solid solution of the FeeMn alloy usually consists of g-austenite (fcc), ε-martensite (hcp) and a0 -
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Please cite this article as: G. Chen et al., Effect of alloying with Al and Cr on the microstructure, damping capacity and high-temperature oxidation behaviors of Fee17Mn damping alloys, Journal of Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2019.153035
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martensite (bct). Several complex phase transformations (g / ε, ε / a0 , g / a0 and ε / g) can occur among these phases [4,5]. Most scholars agree the damping mechanism of FeeMn based alloy is dominated by four factors: (ⅰ) g/ε interface boundaries, (ⅱ) ε-martensite variant boundaries, (ⅲ) stacking faults in g-austenite, and (ⅳ) stacking faults boundaries in ε-martensite. a0 -martensite is usually considered as a harmful phase for damping capacity [6]. Several researchers [5,7] have indicated that the Fee17Mn alloy exhibits its highest damping capacity in the FeeMn binary system because of the superior damping sources. In addition to varying the Mn content, many FeeMn based alloys are prepared by using two or three additional alloying elements. Many researchers studied the effects of additional alloying on the microstructure, damping capacity, corrosion resistance, and mechanical properties of the FeeMn alloys. Wen and Choi [8,9] studied the effects of carbon on the damping capacity of Fee17Mn alloys. They revealed the carbon deteriorated the damping capacity because of segregation. According to Jun et al. [10], alloying Fee23Mn with Co can increase damping capacity by increasing the length of g/ε interface (i.e. increasing the number of ε-martensite plates per unit area). Gavriljuk et al. [11] demonstrated that alloying FeeMn alloys with nitrogen provided a higher level of damping capacity by improving the mobility of interfaces. Huang et al. [12] and Jee et al. [13] pointed out that although the addition of Cr slightly reduced damping capacity, it was also a useful method to improve the corrosion resistance of FeeMn alloy. Besides, the addition of Al into FeeMn alloys can significantly decrease the g / ε transformation while improving corrosion-resistance [14]. It is known that the corrosion resistance of the FeeMn damping alloys is too low for many working conditions. The FeeMn damping alloys exhibit low corrosion resistance because of the co-existence of two or three phases, which easily form electrochemical primary cells. On the other hand, the FeeMn damping alloys can keep a high damping capacity at a wide temperature range from room temperature to 210 C [15]. Thus, it is feasible to apply FeeMn damping alloys in high temperature environment. For example, there is a strong requirement on effectively and accurately suppressing the mechanical vibration and equipment noise to ensure the quiet operation of nuclear power systems [16]. But the in-service temperature of the light water reactor is about 360e415 C, and sometimes the temperature will reach 450e500 C [17]. Therefore, it is necessary to analyze and improve the high-temperature oxidation resistance of the FeeMn damping alloys. In this paper, Fee17Mne5Cr and Fee17Mne5Cre1Al alloys were designed to comparatively study the effects of the addition of Cr and Al on the microstructure, damping capacity, and hightemperature oxidation behaviors of Fee17Mn based damping alloys. And the tests performed at 500 C was to evaluate the oxidation behaviors of materials under the critical operation temperature of nuclear reactors. The effects of alloying the FeeMn damping alloy with Cr and Al were studied by means of X-ray diffraction, SEM, DSC measurement, dynamic mechanical analyzer, and X-ray photoelectron spectroscopy. 2. Experimental methodology 2.1. Materials and solution treatment Three Fee17Mn damping alloys with different Cr and Al contents were prepared (the chemical compositions of these alloys are listed in Table 1). Ingots were prepared by arc-melting industrial pure iron and electrolytic manganese in a ZG-25A vacuum induction melting furnace. Prior to hot forging into 20-mm-thick plates at 1100 C (thermal deformation was not below 70%), the ingots were homogenized at 1150 C for 12 h. Then, the plates were
Table 1 Elemental compositions of the FeeMn damping alloys. Sample codes
Alloys compositions (wt%)
Mn
Cr
M1 M2 M3
Fee17Mn Fee17Mne5Cr Fee17Mne5Cre1Al
17.21 17.13 17.39
4.80 4.53
Al
Fe
1.03
Bal Bal Bal
exposed to a high-temperature solution treatment (at 1000 C for 1 h) with different cooling modes: water cooling (WC), air cooling (AC) and furnace cooling (FC). According to Refs. [18,19], the best solution temperature for FeeMn based damping alloys is 1000 C. The damping capacity of these three materials obtained by different cooling modes were investigated to choose the adequate cooling method. As seen in Fig. 1, damping capacity was significantly improved after solution treatment for the three cooling modes. However, the degree of improvement was as follows: water cooling was the best, air cooling provided a medium improvement, and furnace cooling was the worst. Therefore, higher cooling rates (water cooling) provide higher damping capacity and, accordingly, the alloys by water cooling were chosen for further studies. 2.2. Microstructure characterization Microstructures were observed by field emission scanning electron microscopy (FE-SEM; JSE-5900LV, Japan). Specimens were mechanically polished, electro-polished in a solution of 20% HClO4 þ 80% C2H5OH, and finally etched in a 1.2% K2S2O5 water solution. The solid phase constituents were determined by X-ray diffraction (XRD; Philips X’Pert 1, Netherlands) using Cu-Ka radiation (l ¼ 1.54 Å) in a 2q range from 35 to 85 . According to the scientific literature [20,21], the integrated intensities of ε (101), g (200) and a0 (110) were used to determine the volume fractions of ε, g and a0 in each sample. 2.3. Damping tests Damping capacity was measured by means of a forced subresonance type PL-DMTA Mk II dynamic mechanical analyzer with a size of 50 mm 2 mm 1 mm. The tests were carried out at 30 C with a 1 Hz vibration frequency and strain amplitude ranging from 0 to 1.4 103. A second damping measurement was also performed on the samples to examine their damping stability. Thus, the results of the damping capacity represented the average value of two tests. Besides, before the testing, we have checked the stability by repeating tests three times. 2.4. High-temperature oxidation tests Prior to oxidation, the surfaces of the samples were mechanically polished by using granulometric sandpapers (320, 600, and 1000 mesh). Then, each specimen was cleaned with acetone to ensure a clean and flat surface, and the dimensions of the samples were recorded by a spiral micrometer. Finally, the hightemperature oxidation tests were conducted at 500 C for 35 h in air atmosphere with a furnace. Every 5 h, three samples of each alloy were weighed, and the average weight gained by oxidation was determined intervals of 5 h. Weight gain rate can be described by the weight change per square centimeter of the specimen per minute. Surface and sectional micrographs of the high-temperature oxidized samples were observed by FE-SEM (JSE-5900LV, Japan) and energy-dispersive X-ray spectrometry (EDS). The structure of the oxide scales was analyzed by grazing incidence XRD (XRD;
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Fig. 1. Effects of the solution treatment cooling modes on the damping capacity for the damping alloys: (a) M1, (b) M2, (c) M3.
Philips X’Pert 1, Netherlands) with an angle of 2 , 0.04 scan step size, using Cu-Ka radiation (l ¼ 1.54 Å), and 2q range from 25 to 95 . X-ray photoelectron spectroscopy (XPS; AXIS Ultra DLD, Kratos) was also used to investigate the structure of the oxide scales. 3. Results and discussion
exhibit several lenticular a0 -martensite that is mostly embedded in ε-martensite because of the preferential orientation [22] for the formation of a0 phase. In accordance to the Burgers and the Kurdjumov-Sachs (KeS) orientation relationships [23], small a0 islands are oriented with respect to the ε variants and the g matrix. These orientation relationships are characterized by the parallelism of a close-packed plane and a close-packed direction (i.e.
3.1. Microstructure analysis As shown in Fig. 2, the XRD results indicate that several variants of martensite formed in g-austenite grains after heat treatment. During the solution treatment at 1000 C, high density of vacancies formed in the g-austenite. When the alloys were directly water cooled to room temperature, most of the vacancies were reversed by the high cooling rate. The vacancies provided many favorable sites for the nucleation of martensite during cooling. The intensity of the peaks in the XRD patterns were clearly affected by the contents of Cr and Al, which induced different solid phases. For example, g (111) and a’ (200) diffraction peaks appeared in the pattern for M3 (Fee17Mne5Cre1Al) alloys at 43.563 and 64.442 , respectively. The contents of g and a’ phases in M3 were higher than those in M1 and M2 alloys. Fig. 3 shows the SEM morphologies for the damping alloys after solution treatment. General and detail views are shown for each kind of alloy. As shown in Figs. 3(a), (b), (d), and (e), the microstructure of M1 and M2 alloys shows only ε-martensite and gaustenite plates, because the content of a0 phase is too small to be observed. The M1 alloy seems to have several thin and short ε-martensite embedded in g-austenite. Apparently, Fig. 3(c) and (f)
Fig. 2. XRD patterns of the M1, M2 and M3 damping alloys.
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Fig. 3. SEM morphologies of the Fee17Mn based damping alloys: (a, d) M1, (b, e) M2, (c, f) M3; (aec) general view, (def) detail view.
ð111Þg k ð110Þa’ and ½1120ε k ½111a’). In these alloys, many ε-martensite variants are parallel to each other, and this phenomenon can be explained by the next famous Shoji-Nishiyama (SeN) orientation relationships [24]: ð111Þg k ð0001Þε and ð110Þg k ð2110Þε. Therefore, Al has bigger effects on the microstructure than Cr. To further study the effects of Cr and Al on the microstructure, the volume fractions of g, ε, and a0 phases in the three alloys were calculated based on XRD patterns, and the results are listed in Table 2. These results demonstrate that alloying with only Cr (M2 alloy) has few effects on the microstructure: the volume fraction of ε-martensite decreased by 3.5%, while g-austenite and a0 phase increased by 2.1% and 1.4%, respectively. Furthermore, the volume fraction of ε-martensite in the M3 alloy decreased to 70.9%, while the g-austenite and a0 phase increased by 17.3% and 11.8%, respectively. Therefore, alloying with only Cr decreased the ε-martensite and increased g and a’ phase. Alloying with both, Al and Cr, significantly promoted the change of the phase contents. The effects of alloying elements on the formation of thermalinduced ε-martensite during solution treatment had been studied by Kim at al [25]. Their results showed that the volume fraction of thermal-induced ε-martensite notably decreased by introducing carbon and a consequent reduction of the Ms temperature. Therefore, to confirm a similar effect of Cr and Al in the FeeMn based alloys, the transformation temperatures (Ms, As, and Af) of the M1 to M3 alloys were determined by DSC measurements. As shown in Fig. 4, alloying with only Cr provoked minor decrements on the transformation temperatures. This agrees with the experimental results of Guerrero et al. [26]. Surprisingly, the transformation temperatures of M3 decreased sharply because of the
Table 2 Volume fractions of ε-martensite, g-austenite and a0 -martensite in the M1, M2 and M3 damping alloys. Samples
e-martensite
g-austenite
a0 -martensite
M1 M2 M3
89.1 85.6 70.9
9.4 11.5 17.3
1.5 2.9 11.8
joint addition of Al and Cr. Besides, the decrement of Ms was much larger than that of As, thereby expanding the temperature difference between As and Ms. Kim et al. [25] pointed out the difference between As and Ms is related to the driving force for the transformation. A larger temperature difference between As and Ms means that higher energy needed to overcome the transformation barriers. The driving force is strongly related to the martensitic transformation temperatures. And the martensitic transformation temperatures depend on the compositions of alloys. Thus, the driving force also depends on the compositions of alloys indirectly. Apparently, the driving force has a stronger dependence on Al content than on Cr content. The transformation barriers in FeeMn based alloys mostly result from the stacking fault energy (SFE), because the stacking faults are the nucleation sites for ε-martensite. The SFE of Fee17Mn alloys increased gradually with the increment of Cr content (0-12 wt%) [27,28] but increased sharply with the increment of Al content [29e31]. Therefore, the SFEs in M1, M2, and M3 keep the next relationship: SFEM1
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Fig. 4. MS, AS and AF temperatures measured by DSC method for: (a) M1 alloy, (b) M2 alloy, and (c) M3 alloy.
alloys increased with the increase of amplitude, but the increment was less significant for high amplitudes, especially in the M3 alloy. Alloying with Cr or Al had negative effects on the damping capacity of the FeeMn binary alloys. Alloying with both, Al and Cr, further decreased the damping capacity. For comparison, the damping capacities at strain amplitudes of 3 104 and 1.1 103 are shown in Fig. 5(b). At the low strain amplitude of 3 104, the damping capacity decreased by 21% after alloying with Cr and decreased by 44% after alloying with both Al and Cr. The damping capacity decreased by 8% after alloying with Cr and decreased by about 24% after alloying with both Al and Cr at the high strain amplitude of 1.1 103. Therefore, the effects of alloying with Cr and Al on damping capacity are different at low and high strain amplitudes. To further study the effects of alloying with Cr and Al on the damping capacity of FeeMn alloys, the classic Granatoe-Lücke (G-L) dislocation pinning model was used. According to the G-L model [35], the total damping (Q1), consisting of the strain-amplitudeindependent part Q1 0 and the strain amplitude-dependent part Q1 h , can be expressed as follows:
Q
1
¼ Q 1 0
Q 1 0 ¼
þ
Q 1 h
BLL4C u 36Gb2
Q 1 h ¼ ðC1 = εÞexpð C2 = εÞ
(1)
(2)
(3)
C1 ¼
rFB L3N 6bEL2C
C2 ¼
FB bELC
(4)
where B is the damping constant for dislocation, L is the total length of dislocation line per unit volume, u is the angular frequency, G is the shear modulus, b is the Burger’s vector of dislocations, FB is the binding force between dislocations and weak pinning points, r is the total length of dislocation line per unit volume that participated in the breakaway process, E is the Young modulus, LC is the average dislocation distance between weak pinning points, and LN is the average dislocation distance between the strong pinning points. The unpinning of dislocations is the main factor affecting damping capacity at high amplitudes. By logarithmic operations, Equation (3) can be expressed by the following equation:
C2 þ lnC1 ln Q 1 h *ε ¼ ε
(5)
where C2 and lnC1 are the slope and intercept in the G-L plots, respectively. To obtain a high Q1 value, the value of C1 should be large, while the value of C2 should be small. Fig. 5(c) illustrates the relationship between 1/ε and ln(Q1 h *ε) in the damping alloys. The linear fitting relationship between 1 ln(Q1 is well satisfied for values above the critical value h *ε) and ε (εCr ¼ 4 104), indicating that the damping capacity of the FeeMn samples can be interpreted by the G-L model when strain
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Fig. 5. Damping capacities for M1 to M3 alloys: (a) damping capacities at different strain amplitudes, (b) comparison of damping capacities at strain amplitudes of 3 104 and 1.1 103, (c) ln(Q1*ε) vs 1/ε with a linear fitted results inset, and (d) Stocking faults probability (S.F.P.) in g-austenite and ε-martensite.
amplitude only beyond εcr. The linear fitting coefficients (C1 and C2) are shown in the inset of Fig. 5(c). After alloying with both Al and Cr, the value of C2 increased with a decrease of Lc, indicating alloying with both Al and Cr shortened the average dislocation distance between weak pinning points. On the other hand, LN is the main parameter determining C1. Alloying with both Cr and Al increased the strong pinning points, while the average dislocation distance between the strong pinning points decreased. Alloying with Cr had a big effect on C2, and alloying with Al had a big effect on C1. The reason is that alloying with Cr mainly increases the weak pinning points, but alloying with Al mainly increases the strong pinning points. When the strain amplitude exceeds a critical value, Shockley partial dislocations break away from weak pinning points, rapidly increasing the damping capacity of the alloy. Thus, alloying with Cr has smaller effects on the decrease of damping capacity for high amplitudes above 4 104. Conversely, the damping at low strain amplitudes mainly results from the amount of stacking faults and phase fraction in alloys. The stocking faults probability (S.F.P.) in ε-martensite and g-austenite for M1 to M3 alloys were calculated (see Fig. 5(d)) based on Warren’s theory and peak-shift method. A detailed description of the calculation procedures can be obtained in Refs. [12,36]. The S.F.P. in ε-martensite and g-austenite slightly decreased after alloying with Cr, but significantly decreased after alloying with both Cr and Al. The reason is that SFE increased slightly in alloying with 5 wt% Cr, while sharply increased with the addition of Al. S.P.F. is inversely proportional to SFE. Besides, many researchers agree that the damping capacity increases with the increment of the amount of martensite [37,38]; accordingly, the fraction of ε-martensite decreased after alloying with both Cr and Al (check the phase
fraction in Table 2). This led to the area reduction of g/ε boundaries. 3.3. High-temperature oxidation resistance 3.3.1. Oxidation kinetics The weight gain versus time plots for the three different alloy compositions oxidized in air at 500 C for 0e35 h are shown in Fig. 6(a). Weight gain always increased with the increase of oxidation time. At the early time of the oxidation, weight gain increased rapidly. Then, the increment of weight gain became very small after a transition period. For example, the weight gain plots of M2 and M3 alloys became almost horizontal at the end of the oxidation procedure. The weight gain rates (WGR) at 5 h, 20 h, and 35 h are shown in Fig. 6(b). The weight gain rates of the damping alloys keep the next relationship for a particular oxidation time: WGRM1 > WGRM2 > WGRM3. Furthermore, weight gain rate decreased with increasing oxidation time for all the alloys. Thus, the kinetics plots can be divided into 3 stages: rapid oxidation at the initial stage (Stage A), oxidation at the intermediate stage (Stage B), and slow oxidation at the last stage (Stage C). It is well known that the process of oxidation is the diffusion of metal ions in and out from the bulk in response to the O2 chemical potential, forming sometimes a complex intermediate phase on the surface [39]. Stage A is the period in which the oxides grew rapidly on the samples surface. As time went on, the oxides formed oxide scales that decreased the speed of new oxides formation (i.e. Stage B). During Stage C, the oxide scales became thick and compact, preventing oxygen from diffusing inward. The M1 alloy only showed Stage A (0 he15 h) and the Stage B (15 he35 h), which means a poor high-temperature oxidation
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Fig. 6. Alloys plots for: (a) high-temperature oxidation kinetics at 500 C, and (b) weight gain rates for 5 h, 20 h, and 35 h.
resistance. The Stage A and B for the M2 alloy occurred at the intervals 0 he10 h and 10 he25 h, respectively. These two stages are shorter than those for the M1 alloy. In addition, Stage C appears after oxidizing in the interval from 25 h to 35 h. This might be due to Cr enhancing the potential of the substrate and the protective Cr oxides formed on the surface. The oxidation resistance of the FeeMn alloy was further improved after the addition of Al. The M3 alloy showed shorter Stage A (0 he5 h) and longer Stage B (5 he20 h) and Stage C (20 he35 h) than those of the M1 and M2 alloys. This indicates that the joint addition of Cr and Al is beneficial for the high-temperature oxidation resistance. The relationship between weight gain and time can be fitted by the next exponential function model:
y ¼ kp ½1 expðt = bÞ
(6)
where kp is the oxidation rate constant and b is the time constant. The fitting results are shown in the curves of Fig. 6(a). The rate constants for M1, M2, and M3 alloys are 3.37 104 g cm2 h1, 1.78 104 g cm2 h1, and 1.36 104 g cm2 h1, respectively. These indicate that the high-temperature oxidation resistance was improved by the sole addition of Cr, and further increased by the simultaneous addition of both Cr and Al. 3.3.2. Oxidation scales morphology Fig. 7 shows the microstructure and composition of the damping alloys surface after oxidation at 500 C for 20 h. As shown in Fig. 8(a), the oxide scales in M1 were severely broken and peeled off, which spontaneously occurred during the air cooling from 500 C to room temperature. The detail view of the edge of the broken oxide scales is shown in Fig. 7(b). There are many cracks in the oxide scales and the edges of the broken oxide scales are neat and sharp. This can be explained as follows: the thermal expansion and contraction coefficients of the oxide film and the base layer are different. Thus high thermal stresses generated, breaking the oxide scales [40]. The element contents of the oxide scales (area M1-1) and the peeled regions (area M1-2) are compared by EDS in Fig. 7(c). The oxide scales of the M1 alloys are mainly iron and manganese oxides. The oxygen content of the peeled regions is about 2.5 times lower than that of the oxide scales, showing that the oxygen content decreases from the outside to the inside. Fig. 7(d) is the general view of the M2 alloy samples surface showing a few light color peeled regions. A magnification of the red square in Fig. 7(d) shows a clear view of the oxide scales and peeling (see Fig. 7(e)). The oxide scales revealed an uneven and warped structure. The peeling exhibited high iron content and about 20 at% oxygen, indicating the peeling was a result of the thermal stress. After alloying with Cr, a reduction of the peeling of oxide scales was observed in the surface morphologies.
Furthermore, the peeling exhibited about 5.6 at% Cr, but the surface of the oxide scales had only 0.86 at% Cr. The microstructure of the oxide scales was covered with an oxide scale of small nodules ranging in sizes from 1 mm to 2 mm. Yuan et al. [41] observed a similar microstructure of the oxide scales in Fee25Mne3Cre3Al0.3C-0.01 N austenitic steel during a high temperature oxidation at 700 C for 8 h. This microstructure was attributed to the Cr inside the FeeMn oxides [41]. The microstructure of the M3 alloy sample surface after oxidation is shown in Fig. 8(g) and (h). No peeling occurred, and the surface of the oxide scales were composed of numbers of oxide clusters. EDS results revealed that the alloy oxide was mainly composed of iron and manganese oxides with small amounts of chromium and aluminum oxides. During the oxidation process, the oxidation of the active metal elements (Mn, Cr, and Al) preceded the oxidation of the inert metal element (Fe). According to Wang et al. [42], the M3 alloy studied in the present work had insufficient aluminum content to form continuous alumina scales. After the depletion of aluminum, diffusion of manganese towards the surface layer took place, and manganese oxide was predominantly formed leading to oxide nodules. To confirm the phase composition of the surface oxide, GIXRD analysis was performed on the surface of the damping alloys. The GIXRD results for the alloys oxidized in an air atmosphere at 500 C for 20 h are shown in Fig. 8(a). The phases were identified by matching the patterns with powder crystallographic data (PDF cards: 87e0721 (a-Ferrite), 77e2363 (MnO), 73e1964 (MnFe2O4), and 78e0390 (Mn2O3)). For the M1 alloy, the diffraction peaks corresponded with aFerrite and the oxides: Mn2O3, MnO (halite), and MnFe2O4 (spinel). The pattern of the M1 alloy exhibited the highest intensity for the a-Ferrite (110) peak, illustrating that the surface mainly contains aFerrite. This was caused by the peeling off oxides and the exposure of the matrix. Ferritic layer was likely produced on the surface of the FeeMn damping alloys at high temperatures, because the vapor pressure of Mn is four orders of magnitude higher than that of Fe. The patterns for the M2 and M3 alloys indicated the samples surface mainly contained Mn2O3 and MnFe2O4. A low intensity of aFerrite (110) was also identified. This low intensity emission probably proceeded from too thin oxide scales or small oxide peeled regions. On the other hand, the M2 and M3 alloys exhibited the highest intensity peak (222) for Mn2O3, indicating that the samples surface mainly contained Mn2O3. Because the formation of the Mn2O3 oxide requires very high oxygen partial pressures [43], the Mn2O3 phase was mainly distributed in the outmost part of the oxide scales. XPS was also applied to analyze the chemical composition and state of Mn on the surface of the samples after oxidation (see Fig. 8(b)). The Mn2p3/2 binding energy values for Mn2þ, Mn3þ, and
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Fig. 7. Surface morphologies and composition of the damping alloys surface after oxidation at 500 C for 20 h: (a, b) M1 surface morphologies, (c) EDS results for areas M1-1 and M1-2, (d, e) M2 surface morphologies, (f) EDS results for areas M2-1 and M2-2, (g, h) M3 surface morphologies, and (i) EDS results for Area M3.
Fig. 8. Results for the damping alloys after oxidation at 500 C for 20 h: (a) Grazing incidence XRD (GIXRD) patterns, and (b) Mn 2p3/2 XPS spectra.
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Mn4þ [44] were located at the intervals of 640.7 eV/-640.8 eV, 641.7 eV/-642.3 eV and 643.1 eV/-643.5 eV, respectively. These three valence states of manganese indicate three oxide compounds: Mn2þ reflects MnO, Mn3þ represents Mn2O3, and Mn4þ indicates MnFe2O4. Obviously, the M1 alloy sample surface mainly contained Mn2þ ions, and the amount of Mn2þ ions in M1 was more than that in the M2 or M3 alloys. This indicates the oxide of the M1 sample surface is MnO, which agrees with the results obtained by GIXRD. Fig. 9 shows the cross-sectional SEM back-scattering images and EDS line scan profiles of the damping alloys after oxidation at 500 C in an air atmosphere for 20 h. As shown in Fig. 9(a), no oxide scales were found adhered to the M1 alloy substrate, but some oxide regions appeared in the top region of the substrate. This proves that the oxygen penetrated the voids containing Mn oxide scales [45] and diffused into the substrate through the phase boundaries. Different phases possess different potentials, and therefore phase interfaces are prone to oxidation-reduction reactions [46]. As shown in Fig. 9(d), the oxide regions are rich in Mn and O. The yellow line scan results illustrate the ferritic layer is Mndepleted and about 2 mm wide. As shown in Fig. 9(b), some oxide scales are adhered to the M2 alloy substrate. The oxide scales were rich in Mn and O with contents of Cr and Fe gradually decreasing from the substrate to the surface (observe Fig. 10(e)). The ferritic layer width decreased to about 1.8 mm after alloying with Cr, and no internal oxidation occurred. The Cr might have enhanced the potential of the substrate and improved the potential difference of the phases [47]; therefore, the diffusion of oxygen into the substrate was difficult. The cross-sectional microstructure of the M3 alloy sample after oxidation is shown in Fig. 9(c). Homogenous and compact thin oxide scales were found to be well adhered to the substrate, indicating that their adhesive capacity was improved by adding Al. The ferritic layer of the M3 alloy sample was about 1 mm wide (see Fig. 9(f)). Although Cr and Al generally decrease from the substrate to the oxide scales, a Cr and Al rich region can be observed in the oxide scale near the substrate ferritic layer (check the orange dotted circle in Fig. 9(f)). Accordingly, it can be inferred that Cr and Al
9
Fig. 10. Schematic illustrations of the oxide scales structure: (a) M1, (b) M2, and (c) M3.
oxides assembled in the interface between the oxide scales and the alloy substrate. These interfacial oxides in-between oxide scales and a-ferritic layer are in agreement with the results of Duh et al. [48]. The Cr and Al oxides used the oxygen diffused from the Mn oxide scale and formed a compact layer making the oxygen diffusion into the substrate more difficult. Moreover, less oxygen to oxidize the Mn in the substrate made the Mn diffusion slower and had a negative impact on the growth of the oxide scales. 3.3.3. Oxidation mechanism To illustrate the effects of alloying with Cr and Al on the hightemperature oxidation resistance of the Fee17Mn based alloys, schematic illustrations of the oxide scales of the damping alloys are depicted in Fig. 10. As shown in Fig. 10(a), the outmost oxide scales of the alloy M1 mainly consist of Mn2O3 and MnFe2O4 with several
Fig. 9. Cross-sectional SEM back-scattering images and EDS line scan profiles for the damping alloys after oxidation at 500 C in an air atmosphere for 20 h: (aec) Cross sectional morphologies; and (def) element distribution for the locations shown by yellow lines a, b, and c. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)
Please cite this article as: G. Chen et al., Effect of alloying with Al and Cr on the microstructure, damping capacity and high-temperature oxidation behaviors of Fee17Mn damping alloys, Journal of Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2019.153035
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scattered voids. This was caused by the much higher growth rate of Mn2O3. Nevertheless, the volume ratio between the Mn2O3 scale and the substrate is 0.48, far lower than 1 [45]. The EDS (Fig. 7(c)) and XRD (Fig. 8(a)) results of the outmost oxide layer also showed MnFe2O4. The oxide scales formed by Mn2O3 and MnFe2O4 were easy to peel off during cooling due to the different thermal expansion and contraction coefficients of the oxide scales and the substrate causing high thermal stresses. Although Mn2O3 is a ptype semiconductor [49], the rapid growth and voids in the oxide scales provoked a long initial rapid stage. Accordingly, the amount of the voids in the cracks decreased with the increase of oxidation time. The weight gain rate decreased and the medium stable oxidation stage started (see Stage B in Fig. 6(a)). Under these circumstances, MnO formed at the interface between the oxide scales and the substrate due to the low oxygen partial pressure. Besides, selective oxidation of Mn also took place at the top of the substrate. As shown in Fig. 10(b), the peeling phenomenon of oxide scales was improved by alloying with Cr. The Cr oxides were present all over the oxide scales with a very high content at the interface between the oxide scales and the Mn depleted region (see EDS results in Figs. 7(f) and 9(e)). Cr oxides decreased the voids and cracks in the oxide scales, making more difficult for the oxygen to diffuse inward. On the other hand, the high content of Cr oxides at the interface also hindered the diffusion of oxygen and Mn. Therefore, this caused the earlier start of the slow oxidation stage (see Stage C in Fig. 6(a)) and the improvement of the oxidation resistance. As shown in Fig. 10(c), the oxide scales hardened and peeling was eradicated after alloying with Cr and Al. This is beneficial for the high-temperature oxidation resistance. The Al oxides behaved as the Cr oxides (i.e. they were present all over the oxide scales with a very high content in the interface between the oxide scales and the Mn depleted region). This hardened the inward oxygen diffusion in a short period of time. Because the diffusion process was hindered by oxides, the oxidation process in the Fee17Mn damping alloys followed an exponential function. 4. Conclusions The effects of alloying with Cr and Al on the high-temperature oxidation resistance of Fee17Mn damping alloys were investigated at 500 C. In addition, the effects of the alloying element on the microstructure and damping capacity of such damping alloys were also studied. Based on the results and discussion, it can be concluded that: (1) Alloying with both Cr and Al changed the Ms temperature of the alloys and affected the solid phase compositions. Lower Ms temperatures produced higher g-austenite and ε-martensite phase fractions. Al had a more significant effect on the reduction of the Ms temperature than Cr, because Al sharply increased the stacking faults energy that acted as a barrier for the g / ε phase transformation. (2) Alloying with both Cr and Al decreased damping capacity at low and high amplitudes. This decrement was a result of the reduction of the stacking faults probability and the ε-martensite. At high amplitudes, the pinning of dislocations was the main factor deteriorating damping capacity. While Cr increased the weak pinning points, Al increased the strong pinning points. Therefore, alloying Al decreased the damping capacity more heavily at high amplitudes. (3) Alloying with both Cr and Al significantly improved the hightemperature oxidation resistance of Fee17Mn damping alloys. This alloying procedure also reduced the thermal stresses in the oxide, and consequently the peeling during cooling was reduced. Alloying with both Cr and Al can hinder
the inward oxygen diffusion by reducing voids and forming an oxide layer at the interface between the oxide scales and the substrate.
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Please cite this article as: G. Chen et al., Effect of alloying with Al and Cr on the microstructure, damping capacity and high-temperature oxidation behaviors of Fee17Mn damping alloys, Journal of Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2019.153035