Optical microscopy examinations showed that the as-quenched microstructure of the alloys A(0Cr) through C(6Cr) was single-phase austenite.
Figure 3.1(a) shows a typical optical micrograph of the alloy C(6Cr), indicating single-phase austenite with annealing twins. Transmission electron microscopy examinations indicated that besides the austenite phase, no evidence of carbides could be detected in the as-quenched alloys A(0Cr) through C(6Cr). Figure 3.1(b) is an optical micrograph of the alloy D(9Cr), revealing the presence of some carbides in the austenite matrix and on the grain boundaries. It seems to imply that Cr could be completely dissolved in the austenite matrix at 1373K as Cr ≤ 6 wt.%, and carbides could be formed at this temperature as Cr up to 9 wt.%.
Potentiodynamic polarization curves for the four Fe-Mn-Al-Cr-C alloys in 3.5%
NaCl solution are shown in Figure 3.2. The electrochemical parameters are summarized in Table 3.2. Ecorr of the alloys with different Cr content was varied from -556 mV to -877 mV. Alloy C(6Cr) has the noblest Ecorr (-556mV).
Similarly, with increasing Cr content from 3 to 6 wt%, the Epp was drastically increased from -224mV to -27mV. However, with further increasing the Cr content to 9 wt%, Epp became more negative (-472mV). The results indicate that the alloy C(6Cr) had the highest resistivity to pitting damage. Passivation could
3.2, it is clearly seen that the passive region increased as Cr content increased from 0 to 6 wt.%, and decreased as Cr content up to 9 wt.%. In order to determine the chemical composition and the valence state of element in passive film formed on the alloys in 3.5% NaCl solution, the technique of AES/XPS was undertaken. Figure 3.3 indicates that the depth-concentration profiles for passive film formed on the alloys A(0Cr), C(6Cr) and D(9Cr). The detection of carbon concentration of outmost layer may be due to surface contamination. The O concentration of the alloy A(0Cr) was much smaller than that of the alloys C(6Cr) and D(9Cr) in depth profile. In Figure 3.3(a), it is clear that passive film was not obvious in the alloy without Cr content. However, in the alloys C(6Cr) and D(9Cr), the O concentration at the surface was much higher than that in the matrix, as illustrated in Figures 3.3(b) and (c). The Mn and Fe contents in the outmost surface were much lower than those in the matrix, but Cr and Al contents were reverse tendency. It means that the Cr and Al enrichment was attributed to the preferential dissolution of unstable oxides of Fe and Mn into electrolyte solution, and then replacement by Cr and Al within the oxide layer.
There were broad peaks of Cr and Al at a depth of 0 to 2nm, which corresponded with the peak of O. This implies that the increase of Cr and Al in oxides is likely to be responsible for the improved stability of the passive film.
However, the AES analysis for passive film can not explain the reason why the
Figure 3.1 (a)
Figure 3.1 (b)
Figure 3.1 Optical micrographs of the Fe-30wt.%Mn-7wt.%Al-(6, 9)wt.%
Cr-1wt.%C alloys. (a) 6 wt.%Cr, and (b) 9 wt.%Cr.
Figure 3.2 Potentiodynamic polarization curves for the five Fe-30wt.%Mn- 7wt.%Al-(0, 3, 6, and 9)wt.%Cr-1wt.%C alloys in 3.5% NaCl solution.
Figure 3.3 (a)
Figure 3.3 (b)
Figure 3.3 (c)
Figure 3.3 AES depth profiles for the passive film of the Fe-30wt.%Mn- 7wt.%Al-(0, 6, and 9)wt.%Cr-1wt.%C alloys. (a) 0 wt.%Cr, (b) 6 wt.%Cr, and (c) 9 wt.%Cr.
Table 3.2 The electrochemical parameters from potentiodynamic polarization curves for the Fe-30wt.%Mn-7wt.%Al-(0, 3, 6 and 9)wt.%Cr- 1wt.%C alloys in 3.5% NaCl solution.(*)
Electrochemical Parameters from Polarization Curves Alloy Ecorr（mV） Ecr（mV） Epp（mV） Ip（A/cm2）
A(0Cr) -877 -- -- --
B(3Cr) -712 -588 -224 4.1E-05 C(6Cr) -556 -518 -27 5.75E-06
D(9Cr) -754 -599 -472 1.78E-05
(*)Ecorr, corrosion potential; Ecr, critical potential for active-passive transition;
Epp, pitting potential; Ip, passive current density, minimum value.
Ecorr of the alloy D(9Cr) decreased (more negative), as indicated in Table 3.2. In order to obtain more information, the alloy with 9 wt.% Cr content was examined by TEM. Figure 3.4(a), a bright-field electron micrograph, clearly reveals the presence of carbides in the austenite matrix. Figure 3.4(b) is a selected-area diffraction pattern taken from a carbide. Compared to the previous studies [20], it is clear that the carbide is (Fe,Mn,Cr)7C3 with lattice parameters a
= 1.398nm and c = 0.452nm. Figures 3.4(c) and (d) are two EDS profiles taken from a (Fe,Mn,Cr)7C3 carbide and the austenite matrix nearby the (Fe,Mn,Cr)7C3
carbide, respectively. The quantitative analyses of ten different EDS profiles showed that the chemical composition of the carbide was Fe-0.56wt.%Al- 40.12wt.%Mn-34.39wt.%Cr (EDS with the thick-window detector can only detect the elements with atomic number 11 or above, hence carbon is unable to be examined by this method), and the chemical composition of the austenite matrix nearby the carbide was Fe-10.51wt.%Al-25.48wt.%Mn-2.73wt.%Cr. It is clearly seen that the Cr content in the (Fe,Mn,Cr)7C3 carbide is up to 34.39wt.%, which is much higher than that in the austenite matrix nearby the (Fe,Mn,Cr)7C3
carbide. It is thus expected that the conspicuous decrease of the Cr content in the austenite matrix was ascribed to the formation of Cr-rich (Fe,Mn,Cr)7C3 carbides.
Therefore, it is deduced that the formation of Cr-rich (Fe,Mn,Cr)7C3 carbides would lead to the decrease of Ecorr and Epp of the alloy D(9Cr).
Figure 3.4 (a)
Figure 3.4 (b)
Figure 3.4 (c)
Figure 3.4 (d)
Figure 3.4 Transmission electron micrographs of the Fe-30wt.%Mn-7wt.%Al- 9wt.%Cr-1wt.%C alloy. (a) BF, and (b) a selected-area diffraction pattern taken from a (Fe,Mn,Cr)7C3 carbide in (a). The zone axis of the (Fe,Mn,Cr)7C3 carbide is [1210]. (c) and (d) two typical EDS profiles obtained from the (Fe,Mn,Cr)7C3 carbide and the γ matrix nearby the (Fe,Mn,Cr)7C3 carbide, respectively.
Based on the preceding results, it is clear that no evidence of passivation could be found in the alloy A(0Cr) and the Ecorr of the alloy was -877mV. This result is similar to that examined by other workers in the as-quenched austenitic Fe-(26.0-32.2)wt.%Mn-(8.3-10.0)wt.%Al-(0.85-1.45)wt.%C alloys in 3.5%
NaCl solution, in which they reported that only narrow passive region could be observed and the Ecorr of the alloys was in the range from -789 to -920 mV [5-9,14]. However, an obvious broad passive region could be detected in the present alloys containing Cr. This shows that the Cr addition is indeed beneficial for the corrosion resistance of the austenitic Fe-Mn-Al-C alloys in NaCl solution, which is in agreement with that examined by other workers in the austenitic Fe-(29.5-31.3)wt.%Mn-(8.4-8.9)wt.%Al-(2.6-2.8)wt.%Cr-(0.98-1.06)wt.%C alloys [5,6]. In addition, the Ecorr value for the alloy B(3Cr) was -712 mV, which is comparable to about -720 mV for the as-quenched Fe-29.5wt.%Mn- 8.4wt.%Al-2.6wt.%Cr-1.06wt.%C alloy in 3.5% NaCl solution reported by other workers [6]. In the present study, we have further shown that the 6.0 wt.% Cr addition was completely dissolved in the Fe-30wt.%Mn-7wt.%Al-1.0wt.%C alloy at 1373K. Thus, the Ecorr and Epp values would be pronouncedly increased to -556 mV and -27mV, respectively, in the alloy C(6Cr). However, when the Cr addition was increased up to 9 wt%, the Cr-rich (Fe,Mn,Cr)7C3 carbides could
(Fe,Mn,Cr)7C3 carbides resulted in the drastic decrease of the Ecorr and Epp
values.
3-4 Conclusions
The corrosion potential (Ecorr) of the as-quenched Fe-30wt.%Mn-7wt.%Al- (0, 3, 6 and 9)wt.%Cr-1wt.%C alloys increased from -877 mV to -556 mV as Cr content increased from 0 to 6 wt.%. However, the Ecorr of the alloy with 9 wt.%
Cr dropped to -754 mV due to the formation of (Fe,Mn,Cr)7C3 carbides in the austenitic matrix and on the grain boundaries. The result indicates that the alloy C (6Cr) exhibited the highest corrosion resistance in 3.5% NaCl solution.
Passivation could be observed for all of the present alloys except for the alloy without Cr content.
References
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Chapter 4.
Formation of Hägg carbide in an
Fe-30Mn-10Al-4Cr-0.45C alloy
Formation of Hägg carbide in an Fe-30Mn-10Al-4Cr-0.45C alloy
Abstract
When the present alloy was aged at 550 , Hägg carbides (M℃ 5C2-type carbides) occurred at a/2〈100〉 anti-phase boundaries of the D03 domains. The Hägg carbide has never been observed by previous workers in FeMnAlC and FeMnAlCrC alloy systems. The orientation relationship between Hägg carbide and D03 matrix was determined to be (5_10)M5C2 //(11_0)D03 and (134_) M5C2
//(102_)D03. The orientation relationship between Hägg carbide and bcc-type phase has also never been reported before.
4-1 Introduction
In previous studies [1-6], it is seen that the as-quenched microstructure of the Fe-(28-34)wt.%Mn-(7.8-11)wt.%Al-(0.54-1.3)wt.%C alloys was single-phase austenite (γ). After being aged at 500~750 for moderate times, ℃ fine and coarse (Fe,Mn)3AlC carbides were found to precipitate coherently within the γ matrix and heterogeneously on the γ/γ grain boundaries, respectively. For convenience, the κ′ carbide and κ carbide were used to represent the (Fe,Mn)3AlC carbide formed coherently within the γ matrix and heterogeneously on the γ/γ grain boundaries [2]. After prolonged aging time within this temperature range, the coarse κ carbides grew into adjacent γ grains through a γ → α (ferrite) + β-Mn reaction, a γ → γ0 (carbon-deficient austenite) + κ reaction, a γ → β-Mn + κ reaction or a γ → α + β-Mn + κ reaction [1-5], depending on the chemical composition and aging temperature. In the FeMnAlC alloys with lower carbon content (i.e. 0.4~0.51wt.%C), the as-quenched microstructure was found to be (γ+α) dual phases [7-9]. After being aged at 550~710 , fine κ′ carbides were found to appear within the γ grai℃ ns, and coarse κ carbides as well as β-Mn precipitates could be observed in the α grains and on the α/α grain boundaries [8, 9]. In 1991, the present workers examined the microstructural developments of an Fe-28.6wt.%Mn-10.1wt.%Al-0.46wt.%C alloy [10]. Consequently, it was found that in the as-quenched condition,
extremely fine D03 domains could be observed within the α grains. This is different from that reported by other workers in the duplex FeMnAlC alloys.
When the alloy was aged at temperatures ranging from 450 to 750 , the ℃ ℃ phase transformation sequence occurring within the α grain was found to be α + D03 → α + D03 + κ → α + B2 + κ → α [10].
In order to improve the corrosion resistance and high-temperature oxidation resistance, the Cr has been added to the austenitic or duplex FeMnAlC alloys [11-15]. Based on these results, it can be generally concluded that the addition of Cr does achieve these results. Additionally, the effects of Cr addition on the microstructures of the austenitic FeMnAlC alloys have also been examined by several researchers [16-17]. In the previous study [16], it is seen that when the Fe-30wt.%Mn-9wt.%Al-5wt.%Cr-0.7wt.%C alloy was aged at 550~750 , the ℃ fine κ′ carbides were formed within the γ grain, and a (M7C3 + D03) → (M7C3 + B2) → (M7C3 + α) reaction occurred on the γ/γ grain boundaries. Besides, when the Fe-28.3wt.%Mn-8.7wt.%Al-5.5wt.%Cr-1wt.%C alloy was aged at 800-1250 , a (γ + Cr℃ 7C3) → γ → (γ + (α + B2 + D03)) reaction occurred within the γ grain and on the γ/γ grain boundaries [17]. In contrast to the studies of the austenitic FeMnAlCrC alloys, information concerning the microstructures of the (γ+α) duplex FeMnAlCrC alloys is very deficient. Therefore, the purpose of this
Fe-30wt.%Mn-10wt.%Al-4wt.%Cr-0.45 wt.%C alloy aged at 550 .℃