A novel high-strength high-ductility and high corrosion-resistance FeAlMnC low-density alloy
Abstract
The as-quenched Fe-8.68 wt.% Al-30.5 wt.% Mn-1.85 wt.% C alloy is plasma nitrided at 500C for 8h. The nitrided layer obtained is 40 m-thick and composed predominantly of AlN with a small amount of Fe4N. The resultant surface hardness (1860 Hv), substrate hardness (550 Hv), ductility (33.6%), and corrosion resistance in 3.5% NaCl solution in the present nitrided alloy are far superior to those obtained previously in optimally nitrided high-strength alloy steels, as well as martensitic and precipitation-hardening stainless steels.
2-1 Introduction
In view of excellent combinations of high-strength and high-ductility as well as without needing expensive strategic alloying elements (e.g. Cr, Ni, Mo), the austenitic Fe-Al-Mn-C quarternary alloys have been attracting tremendous attention. Moreover, due to high aluminum content, the density of the alloys is about 13% lower than conventional steels [1]. Previous studies showed that the as-quenched microstructure of Fe-(7.8-10) wt.% Al-(28-30) wt.% Mn-(0.8-1.3) wt.% C alloys was single-phase austenite (γ) [1-5]. An optimal combination of strength and ductility could be obtained for the Fe-Al-Mn-C alloys aged at 550oC for about 16h [2, 3]. Under these aging conditions, a high density of fine (Fe,Mn)3AlC carbides (κ′-carbides) having L′12 (ordered FCC) structure precipitated coherently within γ matrix without any grain boundary precipitates.
After optimal aging, the ultimate tensile strength (UTS), yield strength (YS) and elongation (El) of the Fe-Al-Mn-C alloys could reach 1130-1250 MPa, 1080-1120 MPa and 33-31%, respectively [2,3].
Recently, we investigated the as-quenched microstructure of the Fe-9.8 wt.%
Al -29 wt.% Mn-(1.45-2.05) wt.% C alloys, and found that an extremely high density of nano-sized κ′-carbides was formed within γ matrix by spinodal decomposition during quenching [6]. This is quite different from that observed in the austenitic Fe-Al-Mn-C (C ≦ 1.3 wt. %) alloys, in which fine κ′-carbides
could only be observed in aged alloys. Due to the pre-existing nano-sized κ′-carbides, the aging temperature and time required to attain the optimal combination of strength and ductility are, respectively, much lower and less than those of the previous Fe-Al-Mn-C (C≦1.3 wt.%) alloys. For example, with almost equivalent elongation, the Fe-9 wt. % Al -28 wt.% Mn-1.8 wt.% C alloy aged at 450C for 12h can possess yield strength about 28% higher than that of the optimally aged Fe-Al-Mn-C (C≦1.3 wt.% ) alloys [7].
Although the austenitic Fe-Al-Mn-C alloys could possess excellent combination of strength and ductility, the corrosion resistance of the alloys was insufficient for applications in aggressive environments [4,5]. Plasma nitriding was widely utilized to improve surface hardness and corrosion resistance of metallic materials [8-19]. However, to date, little information concerning the plasma nitriding treatment for Fe-Al-Mn-C alloys has been reported in the literature. The main purpose of this work is to investigate the characteristics of an Fe-8.68 wt.% Al-30.5 wt.% Mn-1.85 wt.% C alloy after plasma nitriding at 500oC for 8h.
2-2 Experimental Procedure
The Fe-8.68 wt. % Al -30.5 wt.% Mn-1.85 wt.% C alloy was prepared in an air induction furnace. After being homogenized at 1150oC for 6h, the ingot was hot-rolled to a 6-mm-thick plate. The plate was subsequently solution heat-treated at 1200oC for 2h and then quenched into room-temperature water.
The specimen was polished using SiC papers to 2400 grit before plasma nitriding. The plasma nitriding process was performed at 500oC for 8h using an atmosphere of 50% N2 and 50% H2 under a pressure of 130Pa. Scanning electron microscopy (SEM) was used to investigate the surface and cross-sectional morphologies of the nitrided alloy by using JEOL 6500 with 15kv voltage.
Transmission electron microscopy examination was performed on a JEOL JEM-2100 transmission electron microscope (TEM) operating at 200kv. TEM specimens were prepared by means of a double-jet electropolisher with an electrolyte of 60% acetic acid, 30% ethanol and 10% perchloric acid. The polishing temperature was kept in the range from -30 oC to -15 oC, and the current density was kept in the range from 3x10-4 to 4x10-4 A/m2. X-ray diffraction (XRD) was carried out using a Bruker D8 with Cu-Kα radiation. The nitrogen concentration and microhardness of the nitrided alloy were determined by using glow discharge spectrometer (GDS) and Vicker’s indenter at 100gf,
respectively. Potentiodynamic polarization curves were measured in 3.5% NaCl solution at 25oC. Electrochemical polarization curves were obtained by using an EG&G Princeton Applied Research Model 273 galvanostat / potentiostat.
Speciemens with an exposed surface area of ~1cm2 were ground with 2400 grit SiC paper and then with 1.5m Al2O3 powder, washed in distilled water and rinsed in acetone prior to passivation. Poteniodynamic polarization curves were obtained at a potential scan rate of 2 mV/s from -1 V to 3V. A saturated calomel electrode (SCE) and a platinum wire were used as reference and auxiliary electrodes, respectively. The specimens for tensile tests were prepared according to ASTM standards. The gauge length, width and thickness of the tensile test specimens are 50 mm, 12.5 mm and 6 mm, respectively. Tensile tests were carried out at room temperature with an Instron 8501 tensile testing machine at a strain rate of 6.710-4 s-1.
2-3 Results and Discussion
Figure 2.1(a) is a TEM (100)κ′ dark-field image and Figure 2.1(b)~(d) are the corresponding diffraction patterns of the as-quenched alloy, revealing that an extremely high density of nano-sized κ′-carbides can be observed within γ matrix and the nano-sized κ′-carbides were formed by spinodal decomposition during quenching [6,7]. By using a LECO 2000 image analyzer, the average size and volume fraction of the κ′-carbides were determined to be about 10 nm and 38%, respectively. A detailed investigation indicated that when the as-quenched alloy was aged at 500oC for 8h, the alloy could possess excellent combination of strength and ductility with the UTS, YS and El being 1402 MPa, 1298 MPa, and 34.5%, respectively. For achieving the effects of aging and nitriding simultaneously, the plasma nitriding was fixed at 500oC for 8h with various processing pressures and gas compositions. The experiments indicated that the working pressure of 130Pa with gas composition of 50% N2 and 50%
H2 could give rise to the best plasma nitriding results. Figure 2.1(e) is a cross-sectional SEM image of the nitrided alloy, showing that the thickness of the nitrided layer is about 40 μm. The grain boundaries of the substrate are clearly revealed by the nital etchant, while the nitrided layer remains intact.
Moreover, the boundary between nitrided layer and substrate is obscure. Figure 2.1(f) shows XRD result for the nitrided alloy, revealing that besides γ
diffraction peaks, diffraction peaks belonging to AlN and Fe4N can also be detected. Both AlN and Fe4N have FCC structure with lattice parameters of 4.06 nm and 3.79 nm, respectively [20,21]. Moreover, the intensity of the AlN diffraction peaks is much higher than that of Fe4N phase, indicating that the nitrided layer is composed predominantly of AlN phase with significantly less amount of Fe4N phase. Furthermore, the XRD peaks are fairly broadened, which may be due to the large amount of nitrogen incorporated in these phases [11-12, 15-19]. Figure 2.2(a) shows the nitrogen concentration as a function of depth, revealing that at the outmost surface, the nitrogen concentration is as high as about 20 wt.% (48 at.%). The nitrogen concentration gradually decreases with increasing depth. Figure 2.2(b) shows the microhardness of the nitrided alloy as a function of depth. The surface microhardness is extremely high (1860 Hv), and gradually decreases with increasing depth until the substrate value of about 550 Hv. Tensile test indicated that UTS, YS, and El of the nitrided alloy were 1388 MPa, 1286 MPa, and 33.6%, respectively, which are comparable to those obtained for the same alloy aged at 500oC for 8h. By slightly tilting the specimen, the fracture and free surfaces could be observed simultaneously, as illustrated in Figure 2.2(c). High density of dimples can be seen within the austenite + κ′-carbides matrix, and no microvoids or microcracks are observed in the vicinity of the interface between nitrided layer
and substrate. Obviously, the substrate remains ductile and the nitrided layer itself is very compact with good adhesion to the substrate.
Potentiodynamic polarization curves for as-quenched and plasma nitrided alloys in 3.5% NaCl solution are shown in Figure 2.3(a). Evidently, for the untreated alloy (curve I), there is no apparent passivation region. The corrosion current density (icorr) and corrosion potential (Ecorr) are 2×10-6 A/cm2 and -790 mV, respectively. However, an obvious passivation region can be observed for the nitrided alloy (curve II), and icorr is evidently reduced by about three orders of magnitude to 6×10-10 A/cm2 and Ecorr is drastically improved to +50 mV.
Moreover, the values of the pitting corrosion current density (ip) and pitting potential (Epit) for the nitrided alloy are 2×10-7 A/cm2 and +2030 mV, respectively. Apparently, plasma nitriding has resulted in a pronounced enhancement in corrosion resistance. Figures 2.3(b) and 2.3(c) are SEM images of the corroded surfaces, indicating that during polarization the grain boundaries and matrix of the untreated alloy were severely attacked, while only a few very small (0.3 m) corrosion pits (as indicated with arrows in Figure 2.3(c)) were formed for the nitrided alloy.
That the nitrided layer of the present nitrided alloy is composed predominantly AlN with a small amount of Fe4N is a remarkable feature. For many industrial applications requiring high strength, high wear resistance and
high corrosion resistance, the nitrided low-Cr (Cr＜1.2 wt.%) alloy steels (e.g.
AISI 4140, 4340 and 5140) and high-Cr (Cr＞12 wt.%) martensitic stainless steels (e.g. AISI 410) as well as precipitation-hardening (PH) stainless steels( e.g. AISI 17-4PH) were widely used. According to extensive previous studies, the optimal nitriding conditions for the low-Cr steels were 520-550 oC for 4-6h [8-10], while those for high-Cr stainless steels were 400-480 oC for 2-20h [11-19]. The nitrided layer formed in these body-centered cubic (BCC) steels is mainly composed of Fe3N (HCP) and Fe4N (FCC), without or with a trace of CrN (FCC) [8-19]. After optimal nitriding treatment, the surface microhardness of the low-Cr alloy steels and high-Cr stainless steels were between 890-940 Hv and 1000-1350 Hv, respectively, which are far lower than 1860 Hv obtained in the present nitrided alloy. The primary reason is that due to AlN formation in the present nitrided alloy, nitrogen concentration near the surface can reach 20 wt.%; whereas the surface nitrogen concentrations of the optimally nitrided low-Cr alloy steels and high-Cr stainless steels were 5.7-10 wt.% and 10-15 wt.%, respectively [17,22-25]. The hardness of the nitrides generally increases with increasing nitrogen concentration. For instance, the hardness of AlN is 25.7GPa [26], which is much higher than that of Fe3N (11.2-12.4 GPa), and Fe4N (8.6-11.2 GPa) [22, 27]. It is worthwhile to emphasize here that the substrate hardness (550Hv) of the present nitrided
alloy is also much higher than 210-400 Hv obtained in the optimally nitrided high-strength alloy steels and stainless steels [9-16, 18]. The reason is that prior to nitriding, these steels need to temper at 15C above the nitriding temperature [28], and then nitrided at the optimal temperature for a long duration. This would deteriorate the substrate hardness drastically [23]. Detailed comparisons of surface hardness and substrate hardness are listed in Table 1.
The most important indicators for evaluating the corrosion resistance of metallic materials are icorr, ip, Ecorr, and Epit; lower current densities and higher potentials indicate better corrosion resistance [8-19]. Table 1 lists the values of icorr, ip, Ecorr, and Epit obtained using the same SCE in 3.5% NaCl solution at room-temperature for the present nitrided alloy, and the previous results for the optimally nitrided low-Cr alloy steels (including 4140, 4340, and 5140), high-Cr 410 martensitic stainless steels, 17-4 precipitation hardening steels (17-4PH) and austenitic stainless steels (including AISI 304, 316). In Table 1, it is obvious that the corrosion resistance of the present nitrided alloy is far superior to that of the low-Cr alloy steels and high-Cr stainless steels. The reasons are described in detail as follows. It is indicated that the low-Cr alloy steels have the lowest surface and substrate hardness. The reason is that prior to nitriding, these steels need to temper at 15C above the nitriding temperature, and then nitrided at the optimal temperature for a long duration. This would
deteriorate the substrate hardness drastically. Moreover, it has shown that plasma nitriding is effective in improving the tribological properties and surface hardness of 410 and 420 martensitic stainless steels under various testing conditions [11-12]. It is also indicated that when nitriding is carried out below a temperature at which CrN forms, the nitrided layer retains its martensite structure but with a larger lattice parameter than the bulk, so that
“expanded martensite” has been proposed in comparison to “expanded austenite”. For this reason, surface hardness of the optimally nitrided martensite stainless steels is slightly lower than that of austenitic stainless steels.
Additionally, 17-4PH stainless steels have been widely utilized in industries.
However, nitrided 17-4PH stainless steel has low surface hardness and poor tribological properties, which could limit its applications in such areas where contact and wear are involved. Some surface modification methods have been carried out for improving the properties of this kind of steel. Recently, plasma nitriding treatments on 17-4PH steel at lower temperature have been developed.
It indicates that these treatments can provide a considerable improvement in the wear resistance of the precipitation-hardening stainless steels without significantly compromising its desirable corrosion-resistant properties.
Consequently, 17-4PH steel also has the excellent combination of surface hardness and corrosion resistance. Finally, it is clearly shown that austenite
stainless steels have the highest surface hardness than other low-Cr and high-Cr steels. It has been explained that when nitriding temperature is sufficiently low, a nitrogen expanded austenite (namely S-phase) can be produced on the surface of the austenite stainless steel. Due to contain a significant amount of nitrogen in the S-phase, the low temperature nitrided austenitic stainless steels can possess not only considerably increased surface hardness and wear resistance, but also much improved corrosion resistance. For example, when nitriding at a temperature around or below 500oC can produce a thick nitrided case on the austenitic stainless steels surface, which indeed largely improve the surface hardness and wear resistance [15-19]. However, corrosion resistance of the austenitic stainless steels is dramatically reduced after nitriding at higher temperature due to the formation of chromium nitride and the depletion of free chromium in the austenite matrix. Consequently, the higher temperature nitrided stainless steel is thus no longer stainless.
Evidently, under the same testing conditions, the icorr and ip of the present alloy are two or three orders of magnitude lower, while the values of Ecorr and Epit are significantly higher than those of the alloy steels and stainless steels, indicating that the present nitrided alloy has far superior corrosion resistance in 3.5 % NaCl solution. Moreover, the size of the surface corrosion pits of the present nitrided alloy is only about 0.3 m (Figure 2.3(c)), which is much
smaller than that (10-200 m) observed in optimally nitrided alloy steels and stainless steels under similar polarization tests [8-9,12-14]. The lower ip value results in smaller corrosion pits [12, 14], which is in good agreement with the experimental results shown in Table 1. Another important criterion for evaluating the pitting resistance is the difference between Epit and Ecorr, namely
△ E=EpitEcorr [29]. In Table 1, the △ E value for optimally nitrided alloy steels and stainless steels is between +270 ~ +1330 mV, while that for the present nitrided alloy is +1980 mV, which again demonstrates the superior characteristics of the present nitrided alloy, presumably due to the high nitrogen concentration at surface [17].
Another feature of the present study is that after etching the boundary between nitrided layer and substrate was obscure (Figure 2.1(e)) and no microvoids or cracks could be detected between nitrided layer and substrate of the fractured surface (Figure 2.2(c)). This is attributed to the fact that both AlN and Fe4N phases have the same FCC crystal structure as the matrix and
-carbides with very similar lattice parameters, which may result in excellent adhesion between nitrided layer and substrate.
2-4 Conclusions
The as-quenched microstructure of the present alloy is ductile phase containing an extremely high density of nano-sized κ′-carbides formed through spinodal decomposition during quenching. The as-quenched alloy is plasma nitrided at 500C for 8h resulting in the effects of aging and nitriding simultaneously. Furthermore, the resultant 40 m-thick nitrided layer is composed predominantly of AlN, the nitrogen concentration at surface is extremely high up to 20 wt.%. Consequently, the surface microhardness (1860 Hv), substrate hardness (550 Hv), ductility (33.6%) and corrosion resistance in 3.5% NaCl solution of the present nitrided alloy are far superior to those obtained previously for the optimally nitrided high-strength alloy steels as well as martensitic and precipitation-hardening stainless steels.
References
[1] G.S. Krivonogov, M.F. Alekseyenko, G.G. Solov’yeva: Fitz. Metal.
Metalloved 39 (1975) 775.
[5] M. Ruscak, T.P. Perng, Corrosion October (1995) 738.
[6] G.D. Tsay, Y.H. Tuan, C.L. Lin, C.G. Chao, T.F. Liu, Mater. Trans. 52 (2011) 521.
[7] K.M. Chang, C.G. Chao, T.F. Liu, Scripta Mater. 63 (2010) 162.
[8] Y. Li, L. Wang, D. Zhang, L. Shen, Applied Surface Science, 256 (2010) 4149.
[9] T. Savisalo, D.B. Lewis, Q. Luo, M. Bolton, P. Hovsepian, Surf. Coat.
Technol. 202 (2008) 1661
[10] Y. Li, L. Wang, D. Zhang, L. Shen, J. Alloys. Compd. 497 (2010) 285.
[11] P. Corengia, G. Ybarra, C. Moina, A. Cabo, E. Broitman, Surf. Coat.
Technol. 187 (2004) 63.
[12] C.X. Li, T. Bell, Corrosion Science 48 (2006) 2036.
[13] R.F. Liu, M.F. Yan, Mater Des 31 (2010) 2355.
[14] R.F. Liu, M.F. Yan, Surf. Coat. Technol. 204 (2010) 2251.
[15] W. Liang, Applied Surface Science 211 (2003) 308.
[16] L. Shen, L. Wang, Y. Wang, C. Wang, Surf. Coat. Technol. 204 (2010) 3222.
[17] C.X. Li, T. Bell, Corrosion Science 46 (2004) 1527.
[18] H.R. Abedi, M. Salehi, Mater Des 32 (2011) 2100.
[19] M. Olzon-Dionysio, S.D. de Souza, R.L.O. Basso, S. de Souza, Surf. Coat.
Technol. 202 (2008) 3607.
[20] S.H. Sheng, R.F. Zhang, S. Veprek, Acta Materia. 56 (2008) 968.
[21] Y. Utsushikawa, K. Niizuma, J. Alloys. Compd. 222 (1995) 188.
[22] H.A. Wriedt, N.A. Gokcen, R.H. Nafziger, Bull. Alloy Phase Diagram 8 (1987)355
[26] J.K. Park, Y.J. Baik, Materials Letters 62 (2008) 2528.
[27] E.A. Ochoa, C.A. Figueroa, F. Alvarez, Surf. Coat. Technol. 200 (2005)
2165.
[28] W.H. Cubberly, V. Masseria, C.W. Kirkpatrick, B. Sanders, Metal
Handbook V.4 Heat Treating, ninth ed., American Society for Metals, Park, Ohio 44073
[29] A. Neville, T. Hodgkiess, Corrosion Science 38 (1996) 927.
Figure 2.1(a)
Figure 2.1(b)
Figure 2.1(c)
Figure 2.1(d)
Figure 2.1(e)
Figure 2.1(f)
Figure 2.1 (a)TEM (100)κ′ dark-field (DF)image of the as-quenched alloy.
(b)~(d)The selected area diffraction patterns (SADPs) taken from the as-quenched alloy (hkl:γ, hkl: κ′-carbide). The zone axis are [001], [011] and [111], respectively. (e) SEM image of the present nitrided alloy (etched in 5% nital). (f) X-ray diffraction pattern for the present nitrided alloy.
-
Figure 2.2(a)
Figure 2.2(b)
Figure 2.2(c)
Figure 2.2 (a)Nitrogen concentration profile measured by GDS of the present nitrided alloy. (b)Hardness profile of the present nitrided alloy. (c) SEM image of the present nitrided alloy after tensile test.
Figure 2.3(a)
Figure 2.3(b)
Figure 2.3(c)
Figure 2.3 (a)Polarization curves for the present untreated and nitrided alloys in 3.5% NaCl solution. (b)-(c) SEM images of the corroded surfaces for the present untreated and nitrided alloys, respectively.
Icorr (A/cm2) Ip (A/cm2) Ecorr (mV) Epit (mV) △E (mV) surface substrate Alloy Steels 8x10-8~4x10-7 4x10-6~9x10-6 -400~-200 +500~+800 +770~+1000 890~940 275~320 410 (MSS) 6x10-8~6x10-7 8x10-5~8x10-4 -220~-30 +50~+600 +270~+630 1150~1204 210~262 17-4PH(SS) 4.1x10-6~9x10-6 9x10-6~1.3x10-5 -208~-207 +700~+715 +907~+923 1160~1167 360~400 304 (SS) 1.46x10-8~1x10-7 4x10-7~2x10-6 -300~-98 +125~+400 +425~+498 1000~1200 220~250 316 (SS) 1x10-7~3.5x10-7 1x10-5~8x10-5 -330~-83.8 +600~+1200 +683.8~+133
0
1350 220 Present Alloy 6x10-10 2x10-7 +50 +2030 +1980 1860 550 Alloy Polarization test results in 3.5% NaCl solution Hardness (Hv)
Table 2.1 Comparisons of polarization test results in 3.5% NaCl solution and hardness of the plasma nitrided Fe-8.68wt.%Al-30.5wt.%Mn-1.85 wt.%C alloy and the optimally plasma nitrided alloy steels as well as various stainless steels.