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Precipitation sequence and its effect on age hardening of alumina-forming austenitic stainless steel
Author’s Accepted Manuscript Precipitation sequence and its effect on age hardening of alumina-forming austenitic stainless steel Joonoh Moon, Tae-Ho Lee, Yoon-Uk Heo, YoungSoo Han, Jun-Yun Kang, Heon-Young Ha, DongWoo Suh www.elsevier.com/locate/msea
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S0921-5093(15)30254-9 http://dx.doi.org/10.1016/j.msea.2015.08.005 MSA32639
To appear in: Materials Science & Engineering A Received date: 3 July 2015 Revised date: 30 July 2015 Accepted date: 1 August 2015 Cite this article as: Joonoh Moon, Tae-Ho Lee, Yoon-Uk Heo, Young-Soo Han, Jun-Yun Kang, Heon-Young Ha and Dong-Woo Suh, Precipitation sequence and its effect on age hardening of alumina-forming austenitic stainless steel, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2015.08.005 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Precipitation sequence and its effect on age hardening of alumina-forming austenitic stainless steel Joonoh Moon a,*, Tae-Ho Lee a, Yoon-Uk Heo b, Young-Soo Han c, Jun-Yun Kang a, Heon-Young Ha a, Dong-Woo Suh b
a
Ferrous Alloy Department, Advanced Metallic Materials Division, Korea Institute of Materials Science, Changwon, Gyeongnam, 642-831, Republic of Korea
b
Graduate Institute of Ferrous Technology, Pohang University of Science and Technology, 77 Cheongam-ro, Nam-gu, Gyeongbuk 790-784, Republic of Korea
c
Neutron Science Division, Korea Atomic Energy Research Institute, Daejeon 305-353, Republic of Korea
* Corresponding author: Joonoh Moon -. E-mail: [email protected] -. Tel. : +82-55-280-3334 -. Fax : +82-55-280-3599
Abstract The precipitation sequence during ageing of Fe-14Cr-20Ni-0.9Nb-2.5Al based aluminaforming austenitic (AFA) steel was explored through a transmission electron microscopy analysis and a small angle neutron scattering experiment. The samples were aged at 700°C for up to 504 h. Particles of NbC, M23C6 and Ni3Al-type L12 were observed in the early stage of ageing. Metastable L12 particles were formed both in grain interior and along grain boundary. M23C6 carbides precipitated along grain boundary accompanied with precipitation of L12 particles. After ageing for longer than 48 h, particles of B2-NiAl and Laves-Fe2Nb were newly formed. We suggest the possibility of phase transition from L12 to B2 with increase in ageing time. Finally, this study examined the change of mechanical properties during ageing through a Gleeble hot tension test and a Vickers hardness test, and then the relationship between precipitation behavior and mechanical properties was carefully investigated and discussed in terms of precipitation behavior.
Key words: Alumina-forming austenitic (AFA) steel, Ageing, Precipitation behavior, Gleeble hot tension test
1. Introduction The efficiency of energy conversion system such as fossil power plant can be improved through the increase in operation temperature, and thus many investigations have focused on the development of advanced alloys compared to those currently used. Recently, alumina-forming austenitic (AFA) stainless steels have received attention due to their excellent oxidation resistance and high creep strength at high temperature during operation of fossil power plant [14]. The addition of 2 to 4 wt% of aluminum (Al) contributes to form an alumina (Al2O3)-based protective film at the surface, which can improve oxidation resistance compared to chromia (Cr2O3) film of conventional Fe-Cr-Ni austenitic stainless steels [1]. That is, Cr2O3 film tends to be weakened by water vapor, leading to an undesired high loss of the materials, while Al2O3 film is much more stable in water vapor condition compared to Cr2O3. In the engineering application concerns, the mechanical performance of AFA steels, particularly the high temperature properties such as creep behavior, is another important issue. One of the most effective methods to improve the creep strength is the precipitation hardening, and thus many researches have investigated the precipitation behaviors and their effects on creep properties in AFA steels during isothermal ageing. Yamamoto et al.[5] studied the effect of alloying additions on creep resistance of AFA steels based on Fe-20Ni-(12~14)Cr-(2.5~4)Al(0.2~3.3)Nb-0.1C and reported that the addition of Nb up to 1wt% and 2.5 to 3 wt% Al
increased creep resistance by the stabilization of nano-sized NbC and B2-NiAl, respectively. However, additions of Nb greater than 1 wt% decreased creep resistance due to the stabilization of Laves-Fe2Nb and the coarsening of NbC [5]. Zhou et al.[6] showed that a quick coarsening of Laves Fe2Nb lead to undesirable creep properties. Also, the optimization of Nb/C ratio contributed to enable the precipitation of stable and fine NbC particle. Zhou et al.[7] reported that the main strengthening media in AFA steel varied with increasing aging temperature. At temperature below 750°C, NbC particle was the major hardening precipitate, but at temperature above 750°C, Laves-Fe2Nb played a role as a dominant precipitate. Meanwhile, the increase of Ni content in AFA steel does not only promote the precipitation of B2-NiAl, but also Ni3Al-type L12 particle which can be used as a strengthener [8]. Yamamoto et al.[8] investigated the effect of aging treatment of AFA steels based on Fe-32Ni-19Cr-2.5Al-3.3Nb and showed the formation of nano-sized L12 particles dispersed in au austenite matrix after creep testing for 500 h at 750°C. In addition, the amount of L12 particles increased with increase in Ti addition which resulted in the precipitation of Ni3(Al,Ti) particles [9]. Despite many previous researches about the precipitation and its effect in AFA steels during ageing treatment, it is somewhat surprising that few studies on the precipitation sequence during ageing are available in the literatures [3, 10]. Investigation about precipitation sequence during ageing is very important to understand the ageing hardening.
Taking all these into account, here we investigate the precipitation sequence and its effect on strengthening during ageing of Fe-14Cr-20Ni-0.9Nb-2.5Al based AFA steel. The changes in the precipitation behavior were carefully analyzed using a scanning electron microscopy (SEM) and a transmission electron microscopy (TEM). In addition, the variation of particle size distribution during ageing was measured by a small angle neutron scattering (SANS) technique. Finally, we evaluated the mechanical property using a Gleeble hot tension test and a Vickers hardness tests, and discussed about correlation between precipitation sequence and ageing hardening behavior.
2. Experimental procedures The chemical composition of AFA steel examined in this investigation is given in Table 1. Ingot was fabricated using a commercial vacuum-induction melting (VIM) furnace. Ingot was homogenized for 2 h at 1200°C, and then hot forged into bar sample of 50 mm in thickness. Fig. 1 shows a schematic schedule for heat treatment of sample after hot forging. The sample was solution-treated for 30 min at 1250°C and then water quenched. Finally, sample was aged at 700°C for up to 504 h. Aged samples underwent a Gleeble hot tensile test at 700°C and a Vickers hardness test under a load of 200 g (0.2 kgf). The microstructures were observed using a SEM (JSM-7001F, JEOL). Precipitation behavior were observed by a TEM (JEM-2100F, JEOL), and the particles were identified by an energy dispersive spectroscopy (EDS) and a selected area diffraction pattern (SADP) analysis. Thin foil specimen for a TEM analysis were fabricated by twin-jet electrolytic polishing at 20 V and 200 mA with a mixed solution of 10 % perchloric acid and 90 % methanol at -20°C. SANS experiments were performed to measure a particle size distribution with increase in ageing time at HANARO in the Korea Atomic Energy Research Institute (KAERI). Neutrons of wavelengths = 7.49 Å with a wavelength spread (full width at half maximum) of 12% were used. Three different sample-to-detector distances (SDD = 1.16 m, 4.7 m and 19.85 m) were
used to cover the overall Q range of 0.0007 Å–1 < Q < 0.55 Å–1. Q is the scattering vector defined as Q = 4πsinθ/λ (where θ is the Bragg angle and λ is the wavelength of the neutron). To obtain quantitative information of the precipitates, the SANS data were fitted assuming that the scatterers were polydisperse spherical precipitates. For a polydisperse system consisting of spherical particles, the macroscopic differential scattering cross section for spheres is given by ∞
2
dσ 4 (Q) = ∆η 2 ∫ π R3 N ( R ) F 2 (Q, R)dR dΩ 3 0
(1) ,
where σ is the cross section, Ω is the solid angle, ∆η is the scattering contrast (η=scattering length density) and R is the and the radius of the spheres, respectively. Here, F(Q,R) is the form factor for the spheres and N(R) is the size distribution function of the spherical particles. Form factor is the Fourier transform of the scattering length density profile depending on the shape of the particles. In model fitting, mathematically defined shape of the particle should be applied. Among a disk, a rod and a sphere, we have to choose one of them. The shape of the particle existing as precipitates in the experimental alloy is near the sphere. So, here we decided to use a sphere for fitting. The model fitting of the real-space size distribution, log-normal size distribution, to the scattering patterns was performed using a non-linear least-squares fitting program. A simple real-space model consists of a set of two or three distributions of the spheres per curve
2 ln( R) − ln( R0 i ) N ( R) = ∑ N 0i exp −0.5 si i
(2) ,
where N0i, R0, and s are the standard scaling factor, and center and width parameters of this distribution type, respectively.
3. Results and discussion
3.1. Microstructure after solution treatment Fig. 2(a) shows a SEM micrograph of the tested alloy after solution treatment. The matrix consisted of austenite, and some annealing twins were observed. Coarse particles precipitated both along grain boundaries and in grain interior, and they were vertically aligned. These coarse particles were identified as NbC with a rock-salt structure through the TEM observation, SAD pattern analysis and EDS analysis, as shown in Fig. 2(b) and (d). Here, we call these coarse NbC as NbC(Ⅰ). Meanwhile, fine particles smaller than 50 nm in Fig. 2(c) were observed in grain interior as well, and they were also identified as NbC particles. We call these fine NbC as NbC(Ⅱ). Moon et al.[11] investigated the precipitation behavior in AFA steel having similar chemical composition with the alloy tested in this study. With considering the phase diagram presented in the literature [11], it is conceivable that coarse NbC(Ⅰ) particles were nucleated during solidification in VIM process, and they were coarsen during solution treatment at 1250° C, as shown in Fig. 2(b). In addition, it is expected that fine NbC(Ⅱ) particles in Fig. 2(c) were
newly precipitated during cooling after solution treatment.
3.2. Precipitation behavior during ageing
Fig. 3 shows SEM images of the samples after solution treatment and after ageing at 700°C for 2 h, 8 h, 24 h, 96 h, and 168 h. While NbC particles were mainly observed at the initial stage of ageing, newly formed particles made up a majority of the precipitates in the matrix after 96 h ageing. Fig. 4 shows TEM micrographs of the samples aged at 700°C for 2 h. Nano-sized NbC particles were observed in Figs. 4(a) and (b), and randomly distributed in the matrix. Theses NbC particles were precipitated not only during cooling after solution treatment, but also during ageing. We call fine NbC precipitated during ageing as NbC(Ⅲ). Dark field image of Fig. 4(d) and SAD pattern analysis of Fig. 4(c) indicate that fine Ni3Al-type L12 particles were precipitated in the matrix, with high fraction. According to the literature [8], L12 particle was precipitated in AFA steel containing high Ni content (about 30 wt%), as a stable phase. However, AFA steel tested in this study contains relatively low Ni content (below 20 wt%). Yamamoto et al.[5] studied the equilibrium phases in composition similar to the present study. They reported that NbC, Fe2Nb Laves, B2-NiAl and M23C6 particles were precipitated as a stable phase, while L12 particle was not observed. Their thermodynamic calculation result also indicated that L12 particle is not stable in AFA steel containing relatively low Ni (less than 20% of Ni in the literature [5]). Therefore, it is somewhat interesting that metastable L12 particles were precipitated with high fraction at the early stage of ageing, as shown in Fig. 4(d). This may be
due to high coherency between L12 particle and austenite matrix [12]. That is, L12 particle and austenite have an excellent coherency due to the similar crystal structure and lattice parameter, as summarized in Table 2 [12-13]. Fig. 5 shows TEM micrographs of the samples aged at 700°C for 8 h, indicating that NbC and L12 particles were also observed. In addition, Figs. 5(c) and (d) show that L12 particles were homogeneously precipitated in the matrix, while nano-sized NbC(Ⅲ) particles were heterogeneously nucleated at dislocations. Meanwhile, Fig. 6 shows TEM micrographs observing particles precipitated along grain boundary after 8 h ageing. These particles were identified as M23C6 by SAD pattern analysis of Fig. 6(c) and (d). For more detailed analysis, EDS mapping along grain boundary was conducted. Fig. 7 shows the EDS mapping results and indicates that Ni and Al were strongly detected in the particles between the particles containing high Cr. From TEM results in Fig.6, it is conceivable that the particles containing high Cr in Fig. 7(b) were M23C6. The particles detected in Fig. 7(c) and (d) might be L12 (Ni3Al) particles. Yamamoto et al. [5] investigated microstructure evolution and creep resistance in composition similar to the present study, and they found grain boundary M23C6 particles after long time ageing of 79 h. Therefore, the precipitation of M23C6 particle at the early stage of ageing in this study is quite interesting. The precipitation of M23C6 particle was accelerated and then this may be closely related to the precipitation of metastable L12 particle. Fig. 8 is indicating schematically the precipitation mechanism at grain boundary. First,
metastable L12 particles were nucleated, and this resulted in enrichment of Cr with Ni depletion around L12 particles. Then, M23C6 particles were able to nucleate in the Cr enriched region. As a result, M23C6 particles precipitated along grain boundary accompanied with precipitation of L12 particles, as shown in Fig. 7 and 8. Meanwhile, grain boundary L12 particles in Fig. 7 were bigger than those in grain interior of Fig. 4 and 5. This is because grain boundary diffusion for particle coarsening is faster than bulk diffusion through matrix [14]. This study measured the particle size distribution using SANS experiments, which are shown in Figs. 9 and 10. Fig. 9 shows the nuclear scattering cross section with increasing scattering vector in both the as-solution treated sample and aged samples at 700°C for 2 h and 8 h. The nuclear scattering cross section of the as-solution treated sample follows a power law with a constant background [15]. The SANS signals in the Q range below 0.06Å-1 were enhanced due to the precipitation of NbC, L12, M23C6 particles. Fig. 10 shows the distribution of volume fraction of particles in the austenite matrix after solution treatment and ageing at 700°C as a function of particle radius. While only NbC(Ⅰ,Ⅱ) particles were detected after solution treatment, L12 and M23C6 particles were additionally detected after ageing treatment, which is well matched with TEM observation results in Figs. 4, 5 and 6. NbC particles show a bimodal distribution, where coarse NbC(Ⅰ) particles were precipitated and coarsen during solution treatment and fine NbC(Ⅱ,Ⅲ) particles were precipitated during cooling after solution
treatment and during ageing. The width of volume fraction peak of fine particles was very narrow in Fig. 10, indicating homogeneous size distribution. In addition, the size of all particles increased with increase in ageing time due to particle coarsening. Next, we carried out TEM observation of aged sample at 700°C for 168 h, in order to identify the needle-shaped particles shown in Fig. 3(e) and (f). Through TEM analysis shown in Fig. 11, needle-shaped particles were identified as B2-NiAl, and Laves-Fe2Nb particles were additionally observed. Laves and B2 phases were precipitated as a stable phase [5].
3.3. Phase transition from L12 to B2 As mentioned above, even if L12 particle is not a stable phase in the tested alloy, it was precipitated at the early stage of ageing due to excellent coherency with austenite matrix. These metastable L12 particles, in nature, disappeared after some ageing time (48 h in this test), and stable B2 particles having similar chemistry and crystal structure with L12 particles were precipitated as shown in Table 2, and finally covered the matrix. The phase transition from L12 to B2 can be understood as following two processes: First, the solute Ni content in the matrix decreased with increase in L12 precipitation, resulting in the suppression of L12 precipitation and the nucleation of B2 particles. The fraction of stable B2 particles naturally increased with increase in ageing time, and they encroached on metastable L12 particles. Next, the coherency
of metastable L12 particle and austenite became difficult to be maintained as coarsening of L12 particle with increase in ageing time, resulting in phase transition from metastable L12 phase to thermodynamically stable B2 phase.
3.4. Change of mechanical properties during ageing Evaluation of mechanical properties was carried out on both solution treated and aged samples at room temperature and 700°C. Fig. 12 shows elevated temperature (700°C) 0.2% offset yield strength (σ0.2) and reduction of area (%RA) as a function of ageing time. Fig. 13 shows the Vickers hardness plotted as a function of ageing time at 700°C. The results evaluated in this study showed that the change of σ0.2 with increase in ageing time at elevated temperature was similar to that of Vickers hardness at room temperature as shown in Figs. 12 and 13. This study divided ageing behavior into three stages with a change of strength. At the first stage, both σ0.2 at 700°C and hardness at RT greatly increased due to the precipitation of fine particles of NbC and L12, as shown in Figs. 4 and 5. The strengthening effect of fine NbC and L12 particles has been widely reported in the literatures [7, 16-18]. Both σ0.2 at 700°C and hardness at RT increased sluggish between 24 h and 48 h ageing time, probably due to the saturation of precipitation and coarsening of NbC and L12 particles. Both σ0.2 at 700°C and hardness at RT increased sharply again after 48 h ageing, which is indicator of the onset of second stage of
ageing hardening. It is because B2 and Laves particles newly formed in the grain interior, as shown in Fig. 11. Finally, hardness shows a peak around 336 h ageing, and then gradually decreased due to over ageing by coarsening of all secondary particles. Meanwhile, according to the increase in strength, %RA value decreased continuously due to the reduction of ductility.
4. Concluding remarks The precipitation behavior and its effect on strengthening during ageing of Fe-14Cr-20Ni0.9Nb-2.5Al based AFA steel were investigated and the following conclusions were drawn. (1) From TEM analysis, the precipitation sequence at 700°C for tested alloy was confirmed as follows: NbC(Ⅰ,Ⅱ) → NbC(Ⅰ,Ⅱ) + NbC(Ⅲ) + metastable L12 (Ni3Al) → NbC(Ⅰ,Ⅱ,Ⅲ) + metastable L12 (Ni3Al) + M23C6 → NbC(Ⅰ,Ⅱ,Ⅲ) + M23C6 + Laves (Fe2Nb) + B2 (NiAl) (2) Metastable L12 particles were precipitated at the early stage of ageing due to excellent coherency with austenite matrix, but they were encroached by stable B2 particles with increase in ageing time. This study suggested the phase transition from L12 to B2 as following two processes: First, the solute Ni content in the matrix decreased with increase in L12 precipitation, resulting in the suppression of L12 precipitation and the nucleation of B2 particles. As a result, L12 particles were encroached by increasing the fraction of stable B2 particles. Next, the coherency of metastable L12 particle and austenite became difficult to be maintained as coarsening of L12 particle with increase in ageing time, resulting in phase transition from metastable L12 phase to thermodynamically stable B2 phase. (3) Interestingly, the precipitation of M23C6 particle was accelerated as compared to the
previous results in the literature, and the precipitation mechanism of intergranular M23C6 was discussed. First, metastable L12 particles were nucleated at grain boundary, resulting in enrichment of Cr with Ni depletion around L12 particles. Subsequently, M23C6 particles were precipitated in the Cr enrichment region. As a result, M23C6 particles precipitated along grain boundary accompanied with precipitation of L12 particles. (4) Tensile behavior at elevated temperature (700°C) was similar to hardness behavior at room temperature, i.e. the ageing hardening behavior was closely related with the precipitation behavior. At the early stage of ageing, strength gradually increased due to the precipitation of fine particles of NbC(Ⅲ) and L12. The strength increase became sluggish between 24 h and 48 h ageing time, due to the saturation of precipitation and coarsening of NbC(Ⅰ,Ⅱ,Ⅲ) and L12 particles. Next, the strength began to increase again after 48 h ageing because B2 and Laves particles newly formed in the grain interior. Then, the strength shows a peak around 336 h ageing, and gradually decreased due to over ageing by coarsening of all particles. Meanwhile, %RA value decreased continuously due to the reduction of ductility, according to the increase in strength.
Acknowledgements This study was supported financially by the Fundamental Research Program of the Korea Institute of Materials Science (KIMS).
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Tables and Figures List Table 1 Chemical compositions of AFA steel examined in this study, in wt% Table 2 Lattice structures of austenite and intermetallic compounds precipitated in AFA steel [12] Fig. 1 Schematic illustration for heat treatment of tested alloy Fig. 2 Microstructure after solution treatment; (a) SEM micrograph, (b, c) TEM micrographs showing coarse and fine NbC particles with the SAD pattern [011]γ //[011]NbC and (d) EDS analysis for NbC particle Fig. 3 SEM images of the samples; (a) after solution treatment, (b~f) aged at 700°C for 2 h, 8h, 24 h, 96 h and 168 h, respectively Fig. 4 TEM micrographs of the sample aged at 700°C for 2 h; (a, b) bright field images of the austenite matrix, (c) dark field image of Ni3Al and (d) SAD pattern analysis Fig. 5 TEM micrographs of the sample aged at 700°C for 8 h; (a, b) bright field images of the austenite matrix, (c) dark field image of fine NbC particles (d) dark field image of Ni3Al, (e) SAD pattern (z=[001]) and (f) a computer-simulated SAD pattern Fig. 6 TEM micrographs showing M23C6 particles precipitated along grain boundary after aging at 700°C for 8 h; (a) bright field image, (b) dark field image, (c and d) SAD patterns Fig. 7 Result of EDS mapping of each element for grain boundary precipitates; (a) bright field image, (b) Cr, (c) Ni and (d) Al Fig. 8 Schematic diagram for nucleation of precipitates at an austenite grain boundary Fig. 9 Nuclear SANS scattering curves of the as solution treated and aged samples at 700°C Fig. 10 Variation in volume distribution with particle size of precipitates with aging time at 700°C; (a) solution treated sample (b) aged sample for 2 h and (c) aged sample for 8 h Fig. 11 TEM micrograph of the aged sample at 700°C for 168 h; (a) bright field image (b) SAD pattern of Laves-Fe2Nb and (c) SAD pattern of B2-NiAl Fig. 12 Results of Gleeble tensile test at elevated temperature Fig. 13 Room temperature Vickers hardness vs. aging time at 700°C
Table 1 Chemical compositions of AFA steel examined in this study, in wt%
AFA
C
Mn
Si
Cr
Ni
Cu
Al
Nb
Mo
Fe
0.11
2.10
2.0
13.5
19.3
2.89
2.59
0.93
2.53
Bal.
Table 2 Lattice structures of austenite and intermetallic compounds precipitated in AFA steel [12,13] Austenite
Ni3Al
NiAl
Lattice parameter (Å)
3.654
3.571
2.848
Structurebericht name
A1
L12
B2
Space group
Fm-3m
Pm-3m
Pm-3m
Fig. 1 Schematic illustration for heat treatment of tested alloy
Fig. 2 Microstructure after solution treatment; (a) SEM micrograph, (b, c) TEM micrographs showing coarse and fine NbC particles with the SAD pattern and (d) EDS analysis for NbC particle
Fig. 3 SEM images of the samples; (a) after solution treatment, (b~f) aged at 700°C for 2 h, 8h, 24 h, 96 h and 168 h, respectively
Fig. 4 TEM micrographs of the sample aged at 700°C for 2 h; (a, b) bright field images of the austenite matrix, (c) SAD pattern analysis and (d) dark field image of Ni3Al
Fig. 5 TEM micrographs of the sample aged at 700°C for 8 h; (a, b) bright field images of the austenite matrix, (c) dark field image of fine NbC particles (d) dark field image of Ni3Al, (e) SAD pattern (z=[001]) and (f) a computer-simulated SAD pattern
Fig. 6 TEM micrographs showing M23C6 particles precipitated along grain boundary after aging at 700°C for 8 h; (a) bright field image, (b) dark field image, (c and d) SAD patterns
Fig. 7 Result of EDS mapping of each element for grain boundary precipitates; (a) bright field image, (b) Cr, (c) Ni and (d) Al
Fig. 8 Schematic diagram for nucleation of precipitates at an austenite grain boundary
Fig. 9 Nuclear SANS scattering curves of the as solution treated and aged samples at 700°C
Fig. 10 Variation in volume distribution with particle size of precipitates with aging time at 700°C; (a) solution treated sample (b) aged sample for 2 h and (c) aged sample for 8 h
Fig. 11 TEM micrograph of the aged sample at 700°C for 168 h; (a) bright field image (b) SAD pattern of Laves-Fe2Nb and (c) SAD pattern of B2-NiAl
Fig. 12 Results of Gleeble tensile test at elevated temperature
Fig. 13 Room temperature Vickers hardness vs. aging time at 700°C
PII: DOI: Reference:
S0921-5093(15)30254-9 http://dx.doi.org/10.1016/j.msea.2015.08.005 MSA32639
To appear in: Materials Science & Engineering A Received date: 3 July 2015 Revised date: 30 July 2015 Accepted date: 1 August 2015 Cite this article as: Joonoh Moon, Tae-Ho Lee, Yoon-Uk Heo, Young-Soo Han, Jun-Yun Kang, Heon-Young Ha and Dong-Woo Suh, Precipitation sequence and its effect on age hardening of alumina-forming austenitic stainless steel, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2015.08.005 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Precipitation sequence and its effect on age hardening of alumina-forming austenitic stainless steel Joonoh Moon a,*, Tae-Ho Lee a, Yoon-Uk Heo b, Young-Soo Han c, Jun-Yun Kang a, Heon-Young Ha a, Dong-Woo Suh b
a
Ferrous Alloy Department, Advanced Metallic Materials Division, Korea Institute of Materials Science, Changwon, Gyeongnam, 642-831, Republic of Korea
b
Graduate Institute of Ferrous Technology, Pohang University of Science and Technology, 77 Cheongam-ro, Nam-gu, Gyeongbuk 790-784, Republic of Korea
c
Neutron Science Division, Korea Atomic Energy Research Institute, Daejeon 305-353, Republic of Korea
* Corresponding author: Joonoh Moon -. E-mail: [email protected] -. Tel. : +82-55-280-3334 -. Fax : +82-55-280-3599
Abstract The precipitation sequence during ageing of Fe-14Cr-20Ni-0.9Nb-2.5Al based aluminaforming austenitic (AFA) steel was explored through a transmission electron microscopy analysis and a small angle neutron scattering experiment. The samples were aged at 700°C for up to 504 h. Particles of NbC, M23C6 and Ni3Al-type L12 were observed in the early stage of ageing. Metastable L12 particles were formed both in grain interior and along grain boundary. M23C6 carbides precipitated along grain boundary accompanied with precipitation of L12 particles. After ageing for longer than 48 h, particles of B2-NiAl and Laves-Fe2Nb were newly formed. We suggest the possibility of phase transition from L12 to B2 with increase in ageing time. Finally, this study examined the change of mechanical properties during ageing through a Gleeble hot tension test and a Vickers hardness test, and then the relationship between precipitation behavior and mechanical properties was carefully investigated and discussed in terms of precipitation behavior.
Key words: Alumina-forming austenitic (AFA) steel, Ageing, Precipitation behavior, Gleeble hot tension test
1. Introduction The efficiency of energy conversion system such as fossil power plant can be improved through the increase in operation temperature, and thus many investigations have focused on the development of advanced alloys compared to those currently used. Recently, alumina-forming austenitic (AFA) stainless steels have received attention due to their excellent oxidation resistance and high creep strength at high temperature during operation of fossil power plant [14]. The addition of 2 to 4 wt% of aluminum (Al) contributes to form an alumina (Al2O3)-based protective film at the surface, which can improve oxidation resistance compared to chromia (Cr2O3) film of conventional Fe-Cr-Ni austenitic stainless steels [1]. That is, Cr2O3 film tends to be weakened by water vapor, leading to an undesired high loss of the materials, while Al2O3 film is much more stable in water vapor condition compared to Cr2O3. In the engineering application concerns, the mechanical performance of AFA steels, particularly the high temperature properties such as creep behavior, is another important issue. One of the most effective methods to improve the creep strength is the precipitation hardening, and thus many researches have investigated the precipitation behaviors and their effects on creep properties in AFA steels during isothermal ageing. Yamamoto et al.[5] studied the effect of alloying additions on creep resistance of AFA steels based on Fe-20Ni-(12~14)Cr-(2.5~4)Al(0.2~3.3)Nb-0.1C and reported that the addition of Nb up to 1wt% and 2.5 to 3 wt% Al
increased creep resistance by the stabilization of nano-sized NbC and B2-NiAl, respectively. However, additions of Nb greater than 1 wt% decreased creep resistance due to the stabilization of Laves-Fe2Nb and the coarsening of NbC [5]. Zhou et al.[6] showed that a quick coarsening of Laves Fe2Nb lead to undesirable creep properties. Also, the optimization of Nb/C ratio contributed to enable the precipitation of stable and fine NbC particle. Zhou et al.[7] reported that the main strengthening media in AFA steel varied with increasing aging temperature. At temperature below 750°C, NbC particle was the major hardening precipitate, but at temperature above 750°C, Laves-Fe2Nb played a role as a dominant precipitate. Meanwhile, the increase of Ni content in AFA steel does not only promote the precipitation of B2-NiAl, but also Ni3Al-type L12 particle which can be used as a strengthener [8]. Yamamoto et al.[8] investigated the effect of aging treatment of AFA steels based on Fe-32Ni-19Cr-2.5Al-3.3Nb and showed the formation of nano-sized L12 particles dispersed in au austenite matrix after creep testing for 500 h at 750°C. In addition, the amount of L12 particles increased with increase in Ti addition which resulted in the precipitation of Ni3(Al,Ti) particles [9]. Despite many previous researches about the precipitation and its effect in AFA steels during ageing treatment, it is somewhat surprising that few studies on the precipitation sequence during ageing are available in the literatures [3, 10]. Investigation about precipitation sequence during ageing is very important to understand the ageing hardening.
Taking all these into account, here we investigate the precipitation sequence and its effect on strengthening during ageing of Fe-14Cr-20Ni-0.9Nb-2.5Al based AFA steel. The changes in the precipitation behavior were carefully analyzed using a scanning electron microscopy (SEM) and a transmission electron microscopy (TEM). In addition, the variation of particle size distribution during ageing was measured by a small angle neutron scattering (SANS) technique. Finally, we evaluated the mechanical property using a Gleeble hot tension test and a Vickers hardness tests, and discussed about correlation between precipitation sequence and ageing hardening behavior.
2. Experimental procedures The chemical composition of AFA steel examined in this investigation is given in Table 1. Ingot was fabricated using a commercial vacuum-induction melting (VIM) furnace. Ingot was homogenized for 2 h at 1200°C, and then hot forged into bar sample of 50 mm in thickness. Fig. 1 shows a schematic schedule for heat treatment of sample after hot forging. The sample was solution-treated for 30 min at 1250°C and then water quenched. Finally, sample was aged at 700°C for up to 504 h. Aged samples underwent a Gleeble hot tensile test at 700°C and a Vickers hardness test under a load of 200 g (0.2 kgf). The microstructures were observed using a SEM (JSM-7001F, JEOL). Precipitation behavior were observed by a TEM (JEM-2100F, JEOL), and the particles were identified by an energy dispersive spectroscopy (EDS) and a selected area diffraction pattern (SADP) analysis. Thin foil specimen for a TEM analysis were fabricated by twin-jet electrolytic polishing at 20 V and 200 mA with a mixed solution of 10 % perchloric acid and 90 % methanol at -20°C. SANS experiments were performed to measure a particle size distribution with increase in ageing time at HANARO in the Korea Atomic Energy Research Institute (KAERI). Neutrons of wavelengths = 7.49 Å with a wavelength spread (full width at half maximum) of 12% were used. Three different sample-to-detector distances (SDD = 1.16 m, 4.7 m and 19.85 m) were
used to cover the overall Q range of 0.0007 Å–1 < Q < 0.55 Å–1. Q is the scattering vector defined as Q = 4πsinθ/λ (where θ is the Bragg angle and λ is the wavelength of the neutron). To obtain quantitative information of the precipitates, the SANS data were fitted assuming that the scatterers were polydisperse spherical precipitates. For a polydisperse system consisting of spherical particles, the macroscopic differential scattering cross section for spheres is given by ∞
2
dσ 4 (Q) = ∆η 2 ∫ π R3 N ( R ) F 2 (Q, R)dR dΩ 3 0
(1) ,
where σ is the cross section, Ω is the solid angle, ∆η is the scattering contrast (η=scattering length density) and R is the and the radius of the spheres, respectively. Here, F(Q,R) is the form factor for the spheres and N(R) is the size distribution function of the spherical particles. Form factor is the Fourier transform of the scattering length density profile depending on the shape of the particles. In model fitting, mathematically defined shape of the particle should be applied. Among a disk, a rod and a sphere, we have to choose one of them. The shape of the particle existing as precipitates in the experimental alloy is near the sphere. So, here we decided to use a sphere for fitting. The model fitting of the real-space size distribution, log-normal size distribution, to the scattering patterns was performed using a non-linear least-squares fitting program. A simple real-space model consists of a set of two or three distributions of the spheres per curve
2 ln( R) − ln( R0 i ) N ( R) = ∑ N 0i exp −0.5 si i
(2) ,
where N0i, R0, and s are the standard scaling factor, and center and width parameters of this distribution type, respectively.
3. Results and discussion
3.1. Microstructure after solution treatment Fig. 2(a) shows a SEM micrograph of the tested alloy after solution treatment. The matrix consisted of austenite, and some annealing twins were observed. Coarse particles precipitated both along grain boundaries and in grain interior, and they were vertically aligned. These coarse particles were identified as NbC with a rock-salt structure through the TEM observation, SAD pattern analysis and EDS analysis, as shown in Fig. 2(b) and (d). Here, we call these coarse NbC as NbC(Ⅰ). Meanwhile, fine particles smaller than 50 nm in Fig. 2(c) were observed in grain interior as well, and they were also identified as NbC particles. We call these fine NbC as NbC(Ⅱ). Moon et al.[11] investigated the precipitation behavior in AFA steel having similar chemical composition with the alloy tested in this study. With considering the phase diagram presented in the literature [11], it is conceivable that coarse NbC(Ⅰ) particles were nucleated during solidification in VIM process, and they were coarsen during solution treatment at 1250° C, as shown in Fig. 2(b). In addition, it is expected that fine NbC(Ⅱ) particles in Fig. 2(c) were
newly precipitated during cooling after solution treatment.
3.2. Precipitation behavior during ageing
Fig. 3 shows SEM images of the samples after solution treatment and after ageing at 700°C for 2 h, 8 h, 24 h, 96 h, and 168 h. While NbC particles were mainly observed at the initial stage of ageing, newly formed particles made up a majority of the precipitates in the matrix after 96 h ageing. Fig. 4 shows TEM micrographs of the samples aged at 700°C for 2 h. Nano-sized NbC particles were observed in Figs. 4(a) and (b), and randomly distributed in the matrix. Theses NbC particles were precipitated not only during cooling after solution treatment, but also during ageing. We call fine NbC precipitated during ageing as NbC(Ⅲ). Dark field image of Fig. 4(d) and SAD pattern analysis of Fig. 4(c) indicate that fine Ni3Al-type L12 particles were precipitated in the matrix, with high fraction. According to the literature [8], L12 particle was precipitated in AFA steel containing high Ni content (about 30 wt%), as a stable phase. However, AFA steel tested in this study contains relatively low Ni content (below 20 wt%). Yamamoto et al.[5] studied the equilibrium phases in composition similar to the present study. They reported that NbC, Fe2Nb Laves, B2-NiAl and M23C6 particles were precipitated as a stable phase, while L12 particle was not observed. Their thermodynamic calculation result also indicated that L12 particle is not stable in AFA steel containing relatively low Ni (less than 20% of Ni in the literature [5]). Therefore, it is somewhat interesting that metastable L12 particles were precipitated with high fraction at the early stage of ageing, as shown in Fig. 4(d). This may be
due to high coherency between L12 particle and austenite matrix [12]. That is, L12 particle and austenite have an excellent coherency due to the similar crystal structure and lattice parameter, as summarized in Table 2 [12-13]. Fig. 5 shows TEM micrographs of the samples aged at 700°C for 8 h, indicating that NbC and L12 particles were also observed. In addition, Figs. 5(c) and (d) show that L12 particles were homogeneously precipitated in the matrix, while nano-sized NbC(Ⅲ) particles were heterogeneously nucleated at dislocations. Meanwhile, Fig. 6 shows TEM micrographs observing particles precipitated along grain boundary after 8 h ageing. These particles were identified as M23C6 by SAD pattern analysis of Fig. 6(c) and (d). For more detailed analysis, EDS mapping along grain boundary was conducted. Fig. 7 shows the EDS mapping results and indicates that Ni and Al were strongly detected in the particles between the particles containing high Cr. From TEM results in Fig.6, it is conceivable that the particles containing high Cr in Fig. 7(b) were M23C6. The particles detected in Fig. 7(c) and (d) might be L12 (Ni3Al) particles. Yamamoto et al. [5] investigated microstructure evolution and creep resistance in composition similar to the present study, and they found grain boundary M23C6 particles after long time ageing of 79 h. Therefore, the precipitation of M23C6 particle at the early stage of ageing in this study is quite interesting. The precipitation of M23C6 particle was accelerated and then this may be closely related to the precipitation of metastable L12 particle. Fig. 8 is indicating schematically the precipitation mechanism at grain boundary. First,
metastable L12 particles were nucleated, and this resulted in enrichment of Cr with Ni depletion around L12 particles. Then, M23C6 particles were able to nucleate in the Cr enriched region. As a result, M23C6 particles precipitated along grain boundary accompanied with precipitation of L12 particles, as shown in Fig. 7 and 8. Meanwhile, grain boundary L12 particles in Fig. 7 were bigger than those in grain interior of Fig. 4 and 5. This is because grain boundary diffusion for particle coarsening is faster than bulk diffusion through matrix [14]. This study measured the particle size distribution using SANS experiments, which are shown in Figs. 9 and 10. Fig. 9 shows the nuclear scattering cross section with increasing scattering vector in both the as-solution treated sample and aged samples at 700°C for 2 h and 8 h. The nuclear scattering cross section of the as-solution treated sample follows a power law with a constant background [15]. The SANS signals in the Q range below 0.06Å-1 were enhanced due to the precipitation of NbC, L12, M23C6 particles. Fig. 10 shows the distribution of volume fraction of particles in the austenite matrix after solution treatment and ageing at 700°C as a function of particle radius. While only NbC(Ⅰ,Ⅱ) particles were detected after solution treatment, L12 and M23C6 particles were additionally detected after ageing treatment, which is well matched with TEM observation results in Figs. 4, 5 and 6. NbC particles show a bimodal distribution, where coarse NbC(Ⅰ) particles were precipitated and coarsen during solution treatment and fine NbC(Ⅱ,Ⅲ) particles were precipitated during cooling after solution
treatment and during ageing. The width of volume fraction peak of fine particles was very narrow in Fig. 10, indicating homogeneous size distribution. In addition, the size of all particles increased with increase in ageing time due to particle coarsening. Next, we carried out TEM observation of aged sample at 700°C for 168 h, in order to identify the needle-shaped particles shown in Fig. 3(e) and (f). Through TEM analysis shown in Fig. 11, needle-shaped particles were identified as B2-NiAl, and Laves-Fe2Nb particles were additionally observed. Laves and B2 phases were precipitated as a stable phase [5].
3.3. Phase transition from L12 to B2 As mentioned above, even if L12 particle is not a stable phase in the tested alloy, it was precipitated at the early stage of ageing due to excellent coherency with austenite matrix. These metastable L12 particles, in nature, disappeared after some ageing time (48 h in this test), and stable B2 particles having similar chemistry and crystal structure with L12 particles were precipitated as shown in Table 2, and finally covered the matrix. The phase transition from L12 to B2 can be understood as following two processes: First, the solute Ni content in the matrix decreased with increase in L12 precipitation, resulting in the suppression of L12 precipitation and the nucleation of B2 particles. The fraction of stable B2 particles naturally increased with increase in ageing time, and they encroached on metastable L12 particles. Next, the coherency
of metastable L12 particle and austenite became difficult to be maintained as coarsening of L12 particle with increase in ageing time, resulting in phase transition from metastable L12 phase to thermodynamically stable B2 phase.
3.4. Change of mechanical properties during ageing Evaluation of mechanical properties was carried out on both solution treated and aged samples at room temperature and 700°C. Fig. 12 shows elevated temperature (700°C) 0.2% offset yield strength (σ0.2) and reduction of area (%RA) as a function of ageing time. Fig. 13 shows the Vickers hardness plotted as a function of ageing time at 700°C. The results evaluated in this study showed that the change of σ0.2 with increase in ageing time at elevated temperature was similar to that of Vickers hardness at room temperature as shown in Figs. 12 and 13. This study divided ageing behavior into three stages with a change of strength. At the first stage, both σ0.2 at 700°C and hardness at RT greatly increased due to the precipitation of fine particles of NbC and L12, as shown in Figs. 4 and 5. The strengthening effect of fine NbC and L12 particles has been widely reported in the literatures [7, 16-18]. Both σ0.2 at 700°C and hardness at RT increased sluggish between 24 h and 48 h ageing time, probably due to the saturation of precipitation and coarsening of NbC and L12 particles. Both σ0.2 at 700°C and hardness at RT increased sharply again after 48 h ageing, which is indicator of the onset of second stage of
ageing hardening. It is because B2 and Laves particles newly formed in the grain interior, as shown in Fig. 11. Finally, hardness shows a peak around 336 h ageing, and then gradually decreased due to over ageing by coarsening of all secondary particles. Meanwhile, according to the increase in strength, %RA value decreased continuously due to the reduction of ductility.
4. Concluding remarks The precipitation behavior and its effect on strengthening during ageing of Fe-14Cr-20Ni0.9Nb-2.5Al based AFA steel were investigated and the following conclusions were drawn. (1) From TEM analysis, the precipitation sequence at 700°C for tested alloy was confirmed as follows: NbC(Ⅰ,Ⅱ) → NbC(Ⅰ,Ⅱ) + NbC(Ⅲ) + metastable L12 (Ni3Al) → NbC(Ⅰ,Ⅱ,Ⅲ) + metastable L12 (Ni3Al) + M23C6 → NbC(Ⅰ,Ⅱ,Ⅲ) + M23C6 + Laves (Fe2Nb) + B2 (NiAl) (2) Metastable L12 particles were precipitated at the early stage of ageing due to excellent coherency with austenite matrix, but they were encroached by stable B2 particles with increase in ageing time. This study suggested the phase transition from L12 to B2 as following two processes: First, the solute Ni content in the matrix decreased with increase in L12 precipitation, resulting in the suppression of L12 precipitation and the nucleation of B2 particles. As a result, L12 particles were encroached by increasing the fraction of stable B2 particles. Next, the coherency of metastable L12 particle and austenite became difficult to be maintained as coarsening of L12 particle with increase in ageing time, resulting in phase transition from metastable L12 phase to thermodynamically stable B2 phase. (3) Interestingly, the precipitation of M23C6 particle was accelerated as compared to the
previous results in the literature, and the precipitation mechanism of intergranular M23C6 was discussed. First, metastable L12 particles were nucleated at grain boundary, resulting in enrichment of Cr with Ni depletion around L12 particles. Subsequently, M23C6 particles were precipitated in the Cr enrichment region. As a result, M23C6 particles precipitated along grain boundary accompanied with precipitation of L12 particles. (4) Tensile behavior at elevated temperature (700°C) was similar to hardness behavior at room temperature, i.e. the ageing hardening behavior was closely related with the precipitation behavior. At the early stage of ageing, strength gradually increased due to the precipitation of fine particles of NbC(Ⅲ) and L12. The strength increase became sluggish between 24 h and 48 h ageing time, due to the saturation of precipitation and coarsening of NbC(Ⅰ,Ⅱ,Ⅲ) and L12 particles. Next, the strength began to increase again after 48 h ageing because B2 and Laves particles newly formed in the grain interior. Then, the strength shows a peak around 336 h ageing, and gradually decreased due to over ageing by coarsening of all particles. Meanwhile, %RA value decreased continuously due to the reduction of ductility, according to the increase in strength.
Acknowledgements This study was supported financially by the Fundamental Research Program of the Korea Institute of Materials Science (KIMS).
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Tables and Figures List Table 1 Chemical compositions of AFA steel examined in this study, in wt% Table 2 Lattice structures of austenite and intermetallic compounds precipitated in AFA steel [12] Fig. 1 Schematic illustration for heat treatment of tested alloy Fig. 2 Microstructure after solution treatment; (a) SEM micrograph, (b, c) TEM micrographs showing coarse and fine NbC particles with the SAD pattern [011]γ //[011]NbC and (d) EDS analysis for NbC particle Fig. 3 SEM images of the samples; (a) after solution treatment, (b~f) aged at 700°C for 2 h, 8h, 24 h, 96 h and 168 h, respectively Fig. 4 TEM micrographs of the sample aged at 700°C for 2 h; (a, b) bright field images of the austenite matrix, (c) dark field image of Ni3Al and (d) SAD pattern analysis Fig. 5 TEM micrographs of the sample aged at 700°C for 8 h; (a, b) bright field images of the austenite matrix, (c) dark field image of fine NbC particles (d) dark field image of Ni3Al, (e) SAD pattern (z=[001]) and (f) a computer-simulated SAD pattern Fig. 6 TEM micrographs showing M23C6 particles precipitated along grain boundary after aging at 700°C for 8 h; (a) bright field image, (b) dark field image, (c and d) SAD patterns Fig. 7 Result of EDS mapping of each element for grain boundary precipitates; (a) bright field image, (b) Cr, (c) Ni and (d) Al Fig. 8 Schematic diagram for nucleation of precipitates at an austenite grain boundary Fig. 9 Nuclear SANS scattering curves of the as solution treated and aged samples at 700°C Fig. 10 Variation in volume distribution with particle size of precipitates with aging time at 700°C; (a) solution treated sample (b) aged sample for 2 h and (c) aged sample for 8 h Fig. 11 TEM micrograph of the aged sample at 700°C for 168 h; (a) bright field image (b) SAD pattern of Laves-Fe2Nb and (c) SAD pattern of B2-NiAl Fig. 12 Results of Gleeble tensile test at elevated temperature Fig. 13 Room temperature Vickers hardness vs. aging time at 700°C
Table 1 Chemical compositions of AFA steel examined in this study, in wt%
AFA
C
Mn
Si
Cr
Ni
Cu
Al
Nb
Mo
Fe
0.11
2.10
2.0
13.5
19.3
2.89
2.59
0.93
2.53
Bal.
Table 2 Lattice structures of austenite and intermetallic compounds precipitated in AFA steel [12,13] Austenite
Ni3Al
NiAl
Lattice parameter (Å)
3.654
3.571
2.848
Structurebericht name
A1
L12
B2
Space group
Fm-3m
Pm-3m
Pm-3m
Fig. 1 Schematic illustration for heat treatment of tested alloy
Fig. 2 Microstructure after solution treatment; (a) SEM micrograph, (b, c) TEM micrographs showing coarse and fine NbC particles with the SAD pattern and (d) EDS analysis for NbC particle
Fig. 3 SEM images of the samples; (a) after solution treatment, (b~f) aged at 700°C for 2 h, 8h, 24 h, 96 h and 168 h, respectively
Fig. 4 TEM micrographs of the sample aged at 700°C for 2 h; (a, b) bright field images of the austenite matrix, (c) SAD pattern analysis and (d) dark field image of Ni3Al
Fig. 5 TEM micrographs of the sample aged at 700°C for 8 h; (a, b) bright field images of the austenite matrix, (c) dark field image of fine NbC particles (d) dark field image of Ni3Al, (e) SAD pattern (z=[001]) and (f) a computer-simulated SAD pattern
Fig. 6 TEM micrographs showing M23C6 particles precipitated along grain boundary after aging at 700°C for 8 h; (a) bright field image, (b) dark field image, (c and d) SAD patterns
Fig. 7 Result of EDS mapping of each element for grain boundary precipitates; (a) bright field image, (b) Cr, (c) Ni and (d) Al
Fig. 8 Schematic diagram for nucleation of precipitates at an austenite grain boundary
Fig. 9 Nuclear SANS scattering curves of the as solution treated and aged samples at 700°C
Fig. 10 Variation in volume distribution with particle size of precipitates with aging time at 700°C; (a) solution treated sample (b) aged sample for 2 h and (c) aged sample for 8 h
Fig. 11 TEM micrograph of the aged sample at 700°C for 168 h; (a) bright field image (b) SAD pattern of Laves-Fe2Nb and (c) SAD pattern of B2-NiAl
Fig. 12 Results of Gleeble tensile test at elevated temperature
Fig. 13 Room temperature Vickers hardness vs. aging time at 700°C