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Mechanical properties of Fe3Al-based alloys with addition of carbon, niobium and titanium
Materials Science and Engineering A 423 (2006) 343–349
Mechanical properties of Fe3Al-based alloys with addition of carbon, niobium and titanium Zheng-Rong Zhang a,∗ , Wen-Xi Liu b a
Department of Energy and Safety Engineering, Yokohama National University, Yokohama 240-8501, Japan b School of Material Science and Engineering, Tianjin University, Tianjin 300072, PR China Received 21 September 2005; received in revised form 7 February 2006; accepted 17 February 2006
Abstract Several Fe3 Al-based iron aluminides with the addition of alloying elements carbon, niobium and titanium were produced by vacuum induction melting (VIM) and hot spinning forging. Analytic techniques including transmission electron microscopy (TEM), scanning electron microscopy (SEM) and X-ray diffraction (XRD) were used in studying the microstructure and fracture manner of these alloys. The results show that due to the addition of alloying elements, the superlattice dislocations tend towards multiple slipping, leaving behind on their slip plane ribbons of square-shaped slip-induced antiphase boundaries. The elongation of Fe3 Al in tension at room temperature increased to about 10% by the addition of suitable alloying elements, the usage of thermo-mechanical processing that has the function of refining grains and substructures, and subsequent annealing. © 2006 Elsevier B.V. All rights reserved. Keywords: Fe3 Al; Iron aluminide; Intermetallics; Tensile ductility; Hot spinning forging
1. Introduction Iron aluminides based on Fe3 Al are promising industrial materials for their excellent oxidation and sulphidation resistance, higher strength and lower density and material cost than those of stainless steels [1–3]. However, their poor ductility at ambient temperatures and an abrupt drop in strength above 600 ◦ C are two limiting factors for potential structural applications. The lower cleavage strength and environmental brittleness were taken as the source of poor values of room temperature elongation which is commonly observed in tests conducted in air [4–9]. The methods that have been used in improving the room temperature ductility include rapid solidification, powder metallurgy techniques, alloying, control of thermo-mechanical processing and heat treatment. Among them, the effects of ternary additions, such as Si, Zr, Mo, Ni, Ti, Nb, V, Cr, Mn and B, on room temperature ductility and strength have been extensively investigated [10–17]. It was found that chromium may increase room temperature ductility while it has almost no effect on strength [18]. Cerium and carbon have also been reported
∗
Corresponding author. Tel.: +81 45 339 3973; fax: +81 45 339 4011. E-mail address: zrzhang [email protected] (Z.-R. Zhang).
0921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2006.02.031
to be beneficial to room temperature ductility of Fe3 Al alloys [19,20]. The objective of the present investigation is to report a change in slip behavior of superlattice dislocations and ductility improvement in Fe3 Al alloy, owing to the addition of alloying elements of carbon, niobium and titanium, and the employment of thermo-mechanical processing. 2. Experimental As good ductility can often be obtained when Fe3 Al is offstoichiometric composition, containing usually 28 at% Al [3], the base composition of alloys used in this study was designed as Fe–28 at% Al with the addition of different amounts of carbon, niobium and titanium. All these alloys were prepared by vacuum induction melting (VIM) and their measured chemical compositions are listed in Table 1. After homogenization at 1000 ◦ C for 6.5 h, the ingots were hot spinning forged at 900–800 ◦ C according to the steps shown in Table 2. During hot spinning forging, the ingot was deformed along three directions as radial deformation (decrease in diameter of the intermediate product), tangential deformation (widening of the intermediate product) and axial deformation (elongation of the intermediate product). Hot spinning forging with the advantages of high frequency pulse forming, small deformation per pulse and extruding along
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Table 1 Chemical composition of Fe3 Al alloys Alloy code
Nominal composition (wt%)
1 2 3
Measured composition (wt%)
Al
C
Nb
Ti
Fe
Al
C
Nb
Ti
Fe
15.82 15.50 15.19
0.06 0.12
0.01 0.01
0.30 0.60
Balance Balance Balance
15.90 15.45 15.36
0.017 0.093 0.11
<0.01 0.027
0.32 0.63
Balance Balance Balance
Table 2 Procedures of spinning forging of Fe3 Al alloys Alloy code ds a (mm) d1 b (mm) d2 (mm) d3 (mm) d4 (mm) df c (mm) 1 2 3
16.50 16.50 16.50 a b c
15 13.5 15
13.5 12.1 13.5
12.1 11.75 12.1
10.75 – 10.75
10.63 11.13 10.54
ds : starting size in diameter. d1 –d4 : inner diameters of forging dies. df : final size in diameter measured at room temperature.
three directions is very suitable to deform brittle intermetallic compounds. All as-forged rods were then kept at 850 ◦ C for 2 h followed by oil cooling. According to the Fe–Al phase diagram [21], ordering heat treatment was performed at 500 ◦ C for 120 h followed by oil cooling to get D03 -phase from B2phase in Fe3 Al. Micro-hardness measurements were performed on Fe3 Al samples at different processing stages, and the average of three readings was plotted for each sample. Specimens with a gauge length of 30 mm and diameter of 5 mm were prepared for tensile tests in air at room temperature with a strain rate of 2.5 × 10−3 s−1 and at 600 ◦ C with a strain rate of 10−4 s−1 , respectively. Microstructure, deformation behavior and fracture mode were characterized on selected tensile specimens using various techniques including transmission electron microscopy (TEM), X-ray diffraction (XRD) and scanning electron microscopy (SEM). For TEM observation, Ø 5 mm × 0.5 mm slice cut off from the tensile specimen was ground to 50–60 m in thickness. Disc with 3 mm diameter punched from the slice was electropolished on twin jet polisher. The electrolyte consists of 15% perchloric acid and 85% ethanol. After ion milling, the disc can be set in JEM-100CX transmission electron microscope for investigation.
be found that after the addition of alloying elements carbon, niobium and titanium, both the yield strength and the tensile strength at room temperature dropped down sharply. This is also consistent with the result of micro-hardness measurements as illustrated in Fig. 1. On the other hand, there are no significant changes in the yield strength and the tensile strength at 600 ◦ C. It should be noted that the room temperature elongation has been doubled for alloyed Fe3 Al even though the absolute value is not high. 3.2. Superlattice dislocations and antiphase boundaries Superlattice rather than ordinary lattice dislocations are the elementary carriers of plastic flow in long range ordered crystals. As illustrated in Fig. 2(a) for a D03 cell of Fe3 Al, where four types of atomic sites on {1 1 0} plane are numbered, the Burgers vector of a perfect dislocation in the D03 structure is a0 1 1 1. Various dissociations into partial superlattice dislocations are possible. The D03 superlattice can be simply constructed by the stacking of two different {1 1 0} planes along 1 1 0 direction (Fig. 2(b)). For the convenience of discussion, two common configurations of superlattice dislocations by the two-fold and four-fold dissociation of a perfect dislocation in the D03 structure are schematized in Fig. 2(c and d) by considering two next near neighboring {1 1 0} planes. There are two different
3. Results 3.1. Tensile properties Tensile tests were carried out at room temperature and at 600 ◦ C, respectively. The results are listed in Table 3. It can
Fig. 1. Micro-hardness of Fe3 Al samples at different processing stages: (A) (䊉) as-cast, (B) () as-hot spinning forged, (C) () after quenching and (D) () after ordering heat treatment.
Table 3 Tensile test results of Fe3 Al alloys performed at room temperature and at 600 ◦ C Alloy code
1 2 3
600 ◦ C
Room temperature Yield strength (MPa)
Tensile strength (MPa)
Elongation (%)
Yield strength (MPa)
Tensile strength (MPa)
Elongation (%)
277.5 160.0 202.5
545.0 412.5 500.0
2.00 4.25 2.00
292.5 297.5 290.0
342.5 335.0 337.5
45.5 28.5 49.5
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Fig. 2. Schematic illustration of (a) four types of atomic sites on {1 1 0} plane in D03 cell of Fe3 Al, (b) two kinds of {1 1 0} stacking planes (atomic sites shown in dark and broken circles), (c) two-fold dissociation of superlattice dislocations and (d) four-fold dissociation of superdislocations. a0 is the lattice constant and Rb and R2 are the displacement vectors corresponding to NNNAPB and NNAPB, respectively.
types of antiphase boundaries (APB) extended between partial superlattice dislocations: a0 /21 1 1 next near neighboring antiphase boundary (NNNAPB) and a0 /41 1 1 nearest neighboring antiphase boundary (NNAPB). Antiphase boundary is a special kind of boundary, which can only be observed by using superlattice reflections rather than by matrix reflections. In this study, only heat-induced a0 /21 0 0 APBs, and slip-induced a0 /21 1 1 APBs, which lie between pairs of partial superlattice dislocations, were observed in tensile specimens tested at room temperature. This is consistent with the results by Yoshimi et al. on APB observation for several Fe3 Al alloys with different chemical compositions and by Yasuda et al. on APB formation in two different Fe3 Al single crystals [22,23]. The slip-induced APBs in alloy 1 are parallel to one another, which reflect the fact that the leading part of the two-fold superlattice dislocations in alloy 1 moves along a straight line. Fig. 3(a) shows two sets of such slip-induced APBs crossing each other. The slip-induced APBs in alloy 2 are square-shaped ribbons (Fig. 3(b)). This indicates that in the presence of alloying elements carbon, niobium and titanium, the leading part of the two-fold superlattice dislocations tend towards multiple slipping, leaving behind on their slip plane ribbons of square shape slip-induced APBs. The heat-induced
APBs in both cases are strip type ribbons. It is noticeable that changes in contrast due to reactions among heat-induced APBs and slip-induced APBs can be observed in the areas where they meet. Fig. 3(c) was taken from the center part of Fig. 3(a) under higher magnification with a slight tilt of specimen, it can be seen that dislocation networks are formed in alloy 1 due to the slip of partial superlattice dislocations. Fig. 3(d) was obtained from the same area of Fig. 3(b) by using different reflections. In alloy 2, the leading part of the two-fold superlattice dislocations tend towards multiple slip and the “walls” of antiphase boundaries are formed. For the specimens tested at 600 ◦ C, dislocations bound together and no slip-induced APBs have been observed. Only heat-induced APBs were observed and the sizes of their corresponding heat-induced antiphase domains (APDs) are smaller than those in specimens after tensile test at room temperature (Fig. 4). 3.3. Fracture mode SEM observation on the topography of Fe3 Al tensile specimens tested at room temperature shows the main fracture feature
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Fig. 3. Superlattice dislocations and APBs in room temperature tensile tested Fe3 Al alloys (a) planar slip in alloy 1, (b) multiple slip in alloy 2, (c) dislocation network in alloy 1 and (d) superlattice dislocations in alloy 2, dark field, g = [8 2¯ 2]* .
Fig. 4. Heat-induced antiphase domains in (a) alloy 1 and (b) alloy 2 after tensile test at 600 ◦ C.
Fig. 5. Transgranular cleavage in Fe3 Al specimens after tensile test at room temperature (a) secondary cracks co-exist with river pattern in alloy 1 and (b) feathering pattern with tearing ridges in alloy 2.
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4. Discussion 4.1. Multiple slip of superlattice dislocations and its contribution to the ductility of Fe3 Al
Fig. 6. Transgranular cleavage with secondary cracks in Fe3 Al specimens after tensile test at 600 ◦ C (a) alloy 1 and (b) alloy 2.
of transgranular cleavage (Fig. 5). River pattern with some secondary cracks distributed on the small facets can be observed in alloy 1. The addition of the alloying elements carbon, niobium and titanium caused the appearance of feathering pattern, which has some tearing ridges in considerable size. At the temperature of 600 ◦ C, the main fracture mode of transgranular cleavage was still kept, as shown in Fig. 6. The secondary cracks in alloy 1 became smaller, but the parallel distribution of many secondary cracks was significant in alloy 2.
Optical microscopy, SEM and XRD observations reveal that the microstructures of all Fe3 Al alloys were basically the same, consisting of single phase of Fe3 Al. No carbide was detected after the addition of alloying elements carbon, niobium and titanium. Instead, a few of flaky secondary phases were formed in alloy 2. They were too small to be inspected by XRD, and were confirmed only by energy dispersive analysis through Xray (EDAX) as Fe3 (Al, Ti). They were also found by TEM observation. As listed in Table 3, after the addition of alloying elements carbon, niobium and titanium, although no positive effect on strengthening Fe3 Al was brought in, the room temperature ductility was indeed effectively improved as the elongation for alloy 2 was more than doubled from 2.0 to 4.25%. This can be explained as the result of a change in slip behavior of superlattice dislocations from planar slip to multiple slip, which will be explained in more details as below. The configuration of typical square-shaped ribbons of APBs is shown in Fig. 7(a and b). This indicates that the multiple slip of superlattice dislocations was dominant in alloy 2. At present, the dissociation and movement of superlattice dislocation in D03 ordered Fe3 Al is still a complex and unclear issue. They are
¯ * , and (c) schematic illustration of the multiple Fig. 7. Square-shaped antiphase boundaries in alloy 2 due to multiple slip (a) bright field, (b) dark field, g = [1 1 1] slip of superlattice dislocation on {1 1 0} planes.
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generally determined by the dislocation core structure and the energy of the APB. The former factor determines the mobility of the superlattice dislocations while the latter one affects the dissociation of superlattice dislocation and interaction between the partial superlattice dislocations [24]. Although the mechanism of the change of antiphase boundary energy in Fe3 Al due to alloying is beyond this work, the experimental result is obvious-only two-fold dissociation of a perfect dislocation in the D03 structure has been observed in present study. And the multiple slip of those superlattice dislocations was frequently observed as dominant slip behavior in alloyed Fe3 Al specimen. In contrast to that of ordinary partial dislocation bounding a stacking fault, the multiple slip of the leading part of the two-fold superlattice dislocations can happen without the constriction of APBs. As illustrated in Fig. 7(c), the leading part of the twofold superlattice dislocations moves contiguously along 1 1 1 direction from one {1 1 0} plane to next perpendicular {1 1 0} planes. The shift at [0 1 0] can be realized by dislocation climb. Because the multiple slip occurs within the perpendicular {1 1 0} slip planes, the projection of APBs on a certain plane would be square shaped. The brittleness of Fe3 Al was often attributed to its lower cleavage strength that connected with the Cottrell mechanism [2]. The formation of a0 [0 1 0] dislocation on (0 0 1) plane, which will trap hydrogen atoms as a stable core of crack and, therefore, is responsible for the lower cleavage strength of materials, was explained as the result of the reaction between two leading parts of two a0 /21 1 1 two-fold superlattice dislocations, which is located on two different {1 1 0} planes. The change in slip behavior of superlattice dislocations from planar slip to multiple slip will certainly make this unlikely to happen. As a result, the cleavage will be restrained. Therefore, the room temperature ductility of alloyed Fe3 Al increased while its elongation was doubled. 4.2. Ductility improvement by thermo-mechanical processing The attempt of improving the ductility of Fe3 Al intermetallic compound was carried out by thermo-mechanical processing on the remaining intermediate product of alloy 3. After homogenization, the rod of alloy 3 was hot spinning forged at 850–800 ◦ C according to the procedure of 13.5 mm → 12.0 mm → 10.5 mm → 9.0 mm → 8.0 mm → 6.5 mm → 4.2 mm → 2.8 mm and wire drawn at room temperature as 2.8 mm → 2.7 mm → 2.6 mm → 2.5 mm → 2.4 mm → 2.3 mm → 2.2 mm → 2.1 mm → 2.0 mm. The wires were then annealed according to the conditions listed in Table 4. Wires with a gauge length of 100 mm and diameter of 2 mm were used in tensile tests in air at room temperature with a strain rate of 2.5 × 10−3 s−1 . The mechanical properties of this alloy at room temperature are displayed in Table 4. It can be concluded that the room temperature elongation of Fe3 Al intermetallic compound achieved at about 10% by the usage of thermo-mechanical processing and subsequent suitable annealing. Elongated fine grains along the direction of wire drown were obtained after thermo-mechanical
Table 4 Mechanical properties of Fe3 Al at room temperature after thermo-mechanical processing and subsequent annealing Specimen code
Diameter of specimen (mm)
Annealing
Tensile strength (MPa)
Elongation (%)
3-1-1 3-1-2
2.00 2.00
600 ◦ C–300 s
1050 1040
10.0 5.0
3-2-1 3-2-2
2.01 2.01
650 ◦ C–300 s
735 755
3.0 5.0
3-3-1 3-3-2
2.01 2.01
700 ◦ C–300 s
675 635
4.0 6.0
3-4-1 3-4-2
1.97 1.99
600 ◦ C–120 s
1020 745
11.0 5.0
3-5-1 3-5-2
2.00 2.00
600 ◦ C–600 s
740 790
10.0 11.0
Fig. 8. Fracture feature of alloy 3 after tensile test at room temperature (a) annealed as 600 ◦ C for 600 s, and (b) amplified image of upper area in (a).
processing. They were slightly recrystallized after annealing at 600 ◦ C and fully recrystallized after annealing at 700 ◦ C. The topography of this alloy after tensile test at room temperature is show in Fig. 8. In addition to intergranular fracture and river pattern of cleavage, some ductile dimples can also be observed. This phenomenon is consistent with the increase in elongation at room temperature. 5. Conclusions (1) The addition of alloying elements of carbon, niobium and titanium into Fe3 Al intermetallic compound results in a change in slip behavior of superlattice dislocations from planar slip to multiple slip, which restrains the cleavage. The dominance of multiple slip of superlattice dislocations is responsible for the room temperature ductility improvement of alloyed Fe3 Al as its elongation being doubled from 2.0 to 4.25%. (2) The dissociation of a perfect dislocation in the D03 structure into two-fold and four-fold configurations of superlattice dislocations and the formation of two different types of antiphase boundaries have been simply schematized by considering two next near neighboring {1 1 0} planes. In this
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study, only heat-induced a0 /21 0 0 APBs and slip-induced a0 /21 1 1 APBs are visible in tensile specimens tested at room temperature. (3) Hot spinning forging, which has the function of refining grains and substructures, combined with subsequent annealing and the selection of suitable alloying elements is a very suitable method to improve the ductility of brittle intermetallic compounds. As a result, the tensile elongation of Fe3 Al at room temperature can be increased to about 10%. Acknowledgement The authors would like to acknowledge Mr. Zheng-Zhong Lei of Division of Research and Development, Tianjin Institute of Metallurgy, for his help on spinning forging performance. References [1] V.K. Sikka, in: L.A. Johnson, D.P. Pope, J.O. Stiegler (Eds.), High Temperature Ordered Intermetallic Alloys IV, Materials Research Society, Pittsburgh, PA, 1991, pp. 907–912. [2] C.G. McKamey, J.H. DeVan, P.F. Tortorelli, V.K. Sikka, J. Mater. Res. 6 (1991) 1779–1805. [3] N.S. Stoloff, Mater. Sci. Eng. A 258 (1998) 1–14. [4] M.G. Mendiratta, S.K. Ehlers, D.K. Chatterjee, H.A. Lipsitt, Metall. Trans. A18 (1987) 283–291. [5] C.G. McKamey, J.A. Horton, C.T. Liu, Scripta Metall. 22 (1988) 1679–1681.
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[6] C.T. Liu, E.H. Lee, C.G. McKamey, Scripta Metall. 23 (1989) 875–880. [7] C.T. Liu, C.G. McKamey, E.H. Lee, Scripta Metall. Mater. 24 (1990) 385–389. [8] R.L. Lynch, L.A. Heldt, W.W. Milligan, Scripta Metall. Mater. 25 (1991) 2147–2151. [9] R.L. Lynch, K.A. Gee, L.A. Heldt, Scripta Metall. Mater. 30 (1994) 945–950. [10] R.G. Bordeau, Development of Iron Aluminides, AFWAL-TR-87-4009, Air Force Wright Aeronautical Laboratories, Wright Patterson Air Force Base, OH, May, 1987. [11] G.E. Fuchs, N.S. Stoloff, Acta Metall. 36 (1988) 1381–1387. [12] C.G. McKamey, J.A. Horton, Metall. Trans. A20 (1989) 751–757. [13] C.G. McKamey, P.J. Maziasz, J.W. Jones, J. Mater. Res. 7 (1992) 2089–2106. [14] S. Yangshan, Z. Zhongua, X. Feng, Y. Xingquan, Mater. Sci. Eng. A 258 (1998) 167–172. [15] Y.D. Huang, W.Y. Yang, Z.Q. Sun, Intermetallics 9 (2001) 119–124. [16] Y. Nishino, T. Tanahashi, Mater. Sci. Eng. A 387–389 (2004) 973–976. [17] D.G. Morris, M.A. Munoz-Morris, C. Baudin, Acta Metall. 52 (2004) 2827–2836. [18] C.G. McKamey, C.T. Liu, Scripta Metall. Mater. 24 (1990) 2119–2122. [19] S. Yangshan, Y. Zhengjun, Z. Zhonghua, H. Haibo, Scripta Metall. Mater. 33 (1995) 811–817. [20] R.G. Baligidad, U. Prakash, A. Radhakrishna, V. Ramakrishna, P.K. Rao, N.B. Ballal, Scripta Mater. 36 (1997) 105–109. [21] S.A. Allen, J.A. Cahn, Acta Metall. 24 (1976) 425. [22] K. Yoshimi, H. Terashima, S. Hanada, Mater. Sci. Eng. A 194 (1995) 53–61. [23] H.Y. Yasuda, K. Nakano, T. Nakajima, M. Ueda, Y. Umakoshi, Acta Metall. 51 (2003) 5101–5112. [24] D. Raabe, J. Keichel, G. Gottstein, Acta Metall. 45 (1997) 2839–2849.
Mechanical properties of Fe3Al-based alloys with addition of carbon, niobium and titanium Zheng-Rong Zhang a,∗ , Wen-Xi Liu b a
Department of Energy and Safety Engineering, Yokohama National University, Yokohama 240-8501, Japan b School of Material Science and Engineering, Tianjin University, Tianjin 300072, PR China Received 21 September 2005; received in revised form 7 February 2006; accepted 17 February 2006
Abstract Several Fe3 Al-based iron aluminides with the addition of alloying elements carbon, niobium and titanium were produced by vacuum induction melting (VIM) and hot spinning forging. Analytic techniques including transmission electron microscopy (TEM), scanning electron microscopy (SEM) and X-ray diffraction (XRD) were used in studying the microstructure and fracture manner of these alloys. The results show that due to the addition of alloying elements, the superlattice dislocations tend towards multiple slipping, leaving behind on their slip plane ribbons of square-shaped slip-induced antiphase boundaries. The elongation of Fe3 Al in tension at room temperature increased to about 10% by the addition of suitable alloying elements, the usage of thermo-mechanical processing that has the function of refining grains and substructures, and subsequent annealing. © 2006 Elsevier B.V. All rights reserved. Keywords: Fe3 Al; Iron aluminide; Intermetallics; Tensile ductility; Hot spinning forging
1. Introduction Iron aluminides based on Fe3 Al are promising industrial materials for their excellent oxidation and sulphidation resistance, higher strength and lower density and material cost than those of stainless steels [1–3]. However, their poor ductility at ambient temperatures and an abrupt drop in strength above 600 ◦ C are two limiting factors for potential structural applications. The lower cleavage strength and environmental brittleness were taken as the source of poor values of room temperature elongation which is commonly observed in tests conducted in air [4–9]. The methods that have been used in improving the room temperature ductility include rapid solidification, powder metallurgy techniques, alloying, control of thermo-mechanical processing and heat treatment. Among them, the effects of ternary additions, such as Si, Zr, Mo, Ni, Ti, Nb, V, Cr, Mn and B, on room temperature ductility and strength have been extensively investigated [10–17]. It was found that chromium may increase room temperature ductility while it has almost no effect on strength [18]. Cerium and carbon have also been reported
∗
Corresponding author. Tel.: +81 45 339 3973; fax: +81 45 339 4011. E-mail address: zrzhang [email protected] (Z.-R. Zhang).
0921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2006.02.031
to be beneficial to room temperature ductility of Fe3 Al alloys [19,20]. The objective of the present investigation is to report a change in slip behavior of superlattice dislocations and ductility improvement in Fe3 Al alloy, owing to the addition of alloying elements of carbon, niobium and titanium, and the employment of thermo-mechanical processing. 2. Experimental As good ductility can often be obtained when Fe3 Al is offstoichiometric composition, containing usually 28 at% Al [3], the base composition of alloys used in this study was designed as Fe–28 at% Al with the addition of different amounts of carbon, niobium and titanium. All these alloys were prepared by vacuum induction melting (VIM) and their measured chemical compositions are listed in Table 1. After homogenization at 1000 ◦ C for 6.5 h, the ingots were hot spinning forged at 900–800 ◦ C according to the steps shown in Table 2. During hot spinning forging, the ingot was deformed along three directions as radial deformation (decrease in diameter of the intermediate product), tangential deformation (widening of the intermediate product) and axial deformation (elongation of the intermediate product). Hot spinning forging with the advantages of high frequency pulse forming, small deformation per pulse and extruding along
344
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Table 1 Chemical composition of Fe3 Al alloys Alloy code
Nominal composition (wt%)
1 2 3
Measured composition (wt%)
Al
C
Nb
Ti
Fe
Al
C
Nb
Ti
Fe
15.82 15.50 15.19
0.06 0.12
0.01 0.01
0.30 0.60
Balance Balance Balance
15.90 15.45 15.36
0.017 0.093 0.11
<0.01 0.027
0.32 0.63
Balance Balance Balance
Table 2 Procedures of spinning forging of Fe3 Al alloys Alloy code ds a (mm) d1 b (mm) d2 (mm) d3 (mm) d4 (mm) df c (mm) 1 2 3
16.50 16.50 16.50 a b c
15 13.5 15
13.5 12.1 13.5
12.1 11.75 12.1
10.75 – 10.75
10.63 11.13 10.54
ds : starting size in diameter. d1 –d4 : inner diameters of forging dies. df : final size in diameter measured at room temperature.
three directions is very suitable to deform brittle intermetallic compounds. All as-forged rods were then kept at 850 ◦ C for 2 h followed by oil cooling. According to the Fe–Al phase diagram [21], ordering heat treatment was performed at 500 ◦ C for 120 h followed by oil cooling to get D03 -phase from B2phase in Fe3 Al. Micro-hardness measurements were performed on Fe3 Al samples at different processing stages, and the average of three readings was plotted for each sample. Specimens with a gauge length of 30 mm and diameter of 5 mm were prepared for tensile tests in air at room temperature with a strain rate of 2.5 × 10−3 s−1 and at 600 ◦ C with a strain rate of 10−4 s−1 , respectively. Microstructure, deformation behavior and fracture mode were characterized on selected tensile specimens using various techniques including transmission electron microscopy (TEM), X-ray diffraction (XRD) and scanning electron microscopy (SEM). For TEM observation, Ø 5 mm × 0.5 mm slice cut off from the tensile specimen was ground to 50–60 m in thickness. Disc with 3 mm diameter punched from the slice was electropolished on twin jet polisher. The electrolyte consists of 15% perchloric acid and 85% ethanol. After ion milling, the disc can be set in JEM-100CX transmission electron microscope for investigation.
be found that after the addition of alloying elements carbon, niobium and titanium, both the yield strength and the tensile strength at room temperature dropped down sharply. This is also consistent with the result of micro-hardness measurements as illustrated in Fig. 1. On the other hand, there are no significant changes in the yield strength and the tensile strength at 600 ◦ C. It should be noted that the room temperature elongation has been doubled for alloyed Fe3 Al even though the absolute value is not high. 3.2. Superlattice dislocations and antiphase boundaries Superlattice rather than ordinary lattice dislocations are the elementary carriers of plastic flow in long range ordered crystals. As illustrated in Fig. 2(a) for a D03 cell of Fe3 Al, where four types of atomic sites on {1 1 0} plane are numbered, the Burgers vector of a perfect dislocation in the D03 structure is a0 1 1 1. Various dissociations into partial superlattice dislocations are possible. The D03 superlattice can be simply constructed by the stacking of two different {1 1 0} planes along 1 1 0 direction (Fig. 2(b)). For the convenience of discussion, two common configurations of superlattice dislocations by the two-fold and four-fold dissociation of a perfect dislocation in the D03 structure are schematized in Fig. 2(c and d) by considering two next near neighboring {1 1 0} planes. There are two different
3. Results 3.1. Tensile properties Tensile tests were carried out at room temperature and at 600 ◦ C, respectively. The results are listed in Table 3. It can
Fig. 1. Micro-hardness of Fe3 Al samples at different processing stages: (A) (䊉) as-cast, (B) () as-hot spinning forged, (C) () after quenching and (D) () after ordering heat treatment.
Table 3 Tensile test results of Fe3 Al alloys performed at room temperature and at 600 ◦ C Alloy code
1 2 3
600 ◦ C
Room temperature Yield strength (MPa)
Tensile strength (MPa)
Elongation (%)
Yield strength (MPa)
Tensile strength (MPa)
Elongation (%)
277.5 160.0 202.5
545.0 412.5 500.0
2.00 4.25 2.00
292.5 297.5 290.0
342.5 335.0 337.5
45.5 28.5 49.5
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Fig. 2. Schematic illustration of (a) four types of atomic sites on {1 1 0} plane in D03 cell of Fe3 Al, (b) two kinds of {1 1 0} stacking planes (atomic sites shown in dark and broken circles), (c) two-fold dissociation of superlattice dislocations and (d) four-fold dissociation of superdislocations. a0 is the lattice constant and Rb and R2 are the displacement vectors corresponding to NNNAPB and NNAPB, respectively.
types of antiphase boundaries (APB) extended between partial superlattice dislocations: a0 /21 1 1 next near neighboring antiphase boundary (NNNAPB) and a0 /41 1 1 nearest neighboring antiphase boundary (NNAPB). Antiphase boundary is a special kind of boundary, which can only be observed by using superlattice reflections rather than by matrix reflections. In this study, only heat-induced a0 /21 0 0 APBs, and slip-induced a0 /21 1 1 APBs, which lie between pairs of partial superlattice dislocations, were observed in tensile specimens tested at room temperature. This is consistent with the results by Yoshimi et al. on APB observation for several Fe3 Al alloys with different chemical compositions and by Yasuda et al. on APB formation in two different Fe3 Al single crystals [22,23]. The slip-induced APBs in alloy 1 are parallel to one another, which reflect the fact that the leading part of the two-fold superlattice dislocations in alloy 1 moves along a straight line. Fig. 3(a) shows two sets of such slip-induced APBs crossing each other. The slip-induced APBs in alloy 2 are square-shaped ribbons (Fig. 3(b)). This indicates that in the presence of alloying elements carbon, niobium and titanium, the leading part of the two-fold superlattice dislocations tend towards multiple slipping, leaving behind on their slip plane ribbons of square shape slip-induced APBs. The heat-induced
APBs in both cases are strip type ribbons. It is noticeable that changes in contrast due to reactions among heat-induced APBs and slip-induced APBs can be observed in the areas where they meet. Fig. 3(c) was taken from the center part of Fig. 3(a) under higher magnification with a slight tilt of specimen, it can be seen that dislocation networks are formed in alloy 1 due to the slip of partial superlattice dislocations. Fig. 3(d) was obtained from the same area of Fig. 3(b) by using different reflections. In alloy 2, the leading part of the two-fold superlattice dislocations tend towards multiple slip and the “walls” of antiphase boundaries are formed. For the specimens tested at 600 ◦ C, dislocations bound together and no slip-induced APBs have been observed. Only heat-induced APBs were observed and the sizes of their corresponding heat-induced antiphase domains (APDs) are smaller than those in specimens after tensile test at room temperature (Fig. 4). 3.3. Fracture mode SEM observation on the topography of Fe3 Al tensile specimens tested at room temperature shows the main fracture feature
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Fig. 3. Superlattice dislocations and APBs in room temperature tensile tested Fe3 Al alloys (a) planar slip in alloy 1, (b) multiple slip in alloy 2, (c) dislocation network in alloy 1 and (d) superlattice dislocations in alloy 2, dark field, g = [8 2¯ 2]* .
Fig. 4. Heat-induced antiphase domains in (a) alloy 1 and (b) alloy 2 after tensile test at 600 ◦ C.
Fig. 5. Transgranular cleavage in Fe3 Al specimens after tensile test at room temperature (a) secondary cracks co-exist with river pattern in alloy 1 and (b) feathering pattern with tearing ridges in alloy 2.
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4. Discussion 4.1. Multiple slip of superlattice dislocations and its contribution to the ductility of Fe3 Al
Fig. 6. Transgranular cleavage with secondary cracks in Fe3 Al specimens after tensile test at 600 ◦ C (a) alloy 1 and (b) alloy 2.
of transgranular cleavage (Fig. 5). River pattern with some secondary cracks distributed on the small facets can be observed in alloy 1. The addition of the alloying elements carbon, niobium and titanium caused the appearance of feathering pattern, which has some tearing ridges in considerable size. At the temperature of 600 ◦ C, the main fracture mode of transgranular cleavage was still kept, as shown in Fig. 6. The secondary cracks in alloy 1 became smaller, but the parallel distribution of many secondary cracks was significant in alloy 2.
Optical microscopy, SEM and XRD observations reveal that the microstructures of all Fe3 Al alloys were basically the same, consisting of single phase of Fe3 Al. No carbide was detected after the addition of alloying elements carbon, niobium and titanium. Instead, a few of flaky secondary phases were formed in alloy 2. They were too small to be inspected by XRD, and were confirmed only by energy dispersive analysis through Xray (EDAX) as Fe3 (Al, Ti). They were also found by TEM observation. As listed in Table 3, after the addition of alloying elements carbon, niobium and titanium, although no positive effect on strengthening Fe3 Al was brought in, the room temperature ductility was indeed effectively improved as the elongation for alloy 2 was more than doubled from 2.0 to 4.25%. This can be explained as the result of a change in slip behavior of superlattice dislocations from planar slip to multiple slip, which will be explained in more details as below. The configuration of typical square-shaped ribbons of APBs is shown in Fig. 7(a and b). This indicates that the multiple slip of superlattice dislocations was dominant in alloy 2. At present, the dissociation and movement of superlattice dislocation in D03 ordered Fe3 Al is still a complex and unclear issue. They are
¯ * , and (c) schematic illustration of the multiple Fig. 7. Square-shaped antiphase boundaries in alloy 2 due to multiple slip (a) bright field, (b) dark field, g = [1 1 1] slip of superlattice dislocation on {1 1 0} planes.
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generally determined by the dislocation core structure and the energy of the APB. The former factor determines the mobility of the superlattice dislocations while the latter one affects the dissociation of superlattice dislocation and interaction between the partial superlattice dislocations [24]. Although the mechanism of the change of antiphase boundary energy in Fe3 Al due to alloying is beyond this work, the experimental result is obvious-only two-fold dissociation of a perfect dislocation in the D03 structure has been observed in present study. And the multiple slip of those superlattice dislocations was frequently observed as dominant slip behavior in alloyed Fe3 Al specimen. In contrast to that of ordinary partial dislocation bounding a stacking fault, the multiple slip of the leading part of the two-fold superlattice dislocations can happen without the constriction of APBs. As illustrated in Fig. 7(c), the leading part of the twofold superlattice dislocations moves contiguously along 1 1 1 direction from one {1 1 0} plane to next perpendicular {1 1 0} planes. The shift at [0 1 0] can be realized by dislocation climb. Because the multiple slip occurs within the perpendicular {1 1 0} slip planes, the projection of APBs on a certain plane would be square shaped. The brittleness of Fe3 Al was often attributed to its lower cleavage strength that connected with the Cottrell mechanism [2]. The formation of a0 [0 1 0] dislocation on (0 0 1) plane, which will trap hydrogen atoms as a stable core of crack and, therefore, is responsible for the lower cleavage strength of materials, was explained as the result of the reaction between two leading parts of two a0 /21 1 1 two-fold superlattice dislocations, which is located on two different {1 1 0} planes. The change in slip behavior of superlattice dislocations from planar slip to multiple slip will certainly make this unlikely to happen. As a result, the cleavage will be restrained. Therefore, the room temperature ductility of alloyed Fe3 Al increased while its elongation was doubled. 4.2. Ductility improvement by thermo-mechanical processing The attempt of improving the ductility of Fe3 Al intermetallic compound was carried out by thermo-mechanical processing on the remaining intermediate product of alloy 3. After homogenization, the rod of alloy 3 was hot spinning forged at 850–800 ◦ C according to the procedure of 13.5 mm → 12.0 mm → 10.5 mm → 9.0 mm → 8.0 mm → 6.5 mm → 4.2 mm → 2.8 mm and wire drawn at room temperature as 2.8 mm → 2.7 mm → 2.6 mm → 2.5 mm → 2.4 mm → 2.3 mm → 2.2 mm → 2.1 mm → 2.0 mm. The wires were then annealed according to the conditions listed in Table 4. Wires with a gauge length of 100 mm and diameter of 2 mm were used in tensile tests in air at room temperature with a strain rate of 2.5 × 10−3 s−1 . The mechanical properties of this alloy at room temperature are displayed in Table 4. It can be concluded that the room temperature elongation of Fe3 Al intermetallic compound achieved at about 10% by the usage of thermo-mechanical processing and subsequent suitable annealing. Elongated fine grains along the direction of wire drown were obtained after thermo-mechanical
Table 4 Mechanical properties of Fe3 Al at room temperature after thermo-mechanical processing and subsequent annealing Specimen code
Diameter of specimen (mm)
Annealing
Tensile strength (MPa)
Elongation (%)
3-1-1 3-1-2
2.00 2.00
600 ◦ C–300 s
1050 1040
10.0 5.0
3-2-1 3-2-2
2.01 2.01
650 ◦ C–300 s
735 755
3.0 5.0
3-3-1 3-3-2
2.01 2.01
700 ◦ C–300 s
675 635
4.0 6.0
3-4-1 3-4-2
1.97 1.99
600 ◦ C–120 s
1020 745
11.0 5.0
3-5-1 3-5-2
2.00 2.00
600 ◦ C–600 s
740 790
10.0 11.0
Fig. 8. Fracture feature of alloy 3 after tensile test at room temperature (a) annealed as 600 ◦ C for 600 s, and (b) amplified image of upper area in (a).
processing. They were slightly recrystallized after annealing at 600 ◦ C and fully recrystallized after annealing at 700 ◦ C. The topography of this alloy after tensile test at room temperature is show in Fig. 8. In addition to intergranular fracture and river pattern of cleavage, some ductile dimples can also be observed. This phenomenon is consistent with the increase in elongation at room temperature. 5. Conclusions (1) The addition of alloying elements of carbon, niobium and titanium into Fe3 Al intermetallic compound results in a change in slip behavior of superlattice dislocations from planar slip to multiple slip, which restrains the cleavage. The dominance of multiple slip of superlattice dislocations is responsible for the room temperature ductility improvement of alloyed Fe3 Al as its elongation being doubled from 2.0 to 4.25%. (2) The dissociation of a perfect dislocation in the D03 structure into two-fold and four-fold configurations of superlattice dislocations and the formation of two different types of antiphase boundaries have been simply schematized by considering two next near neighboring {1 1 0} planes. In this
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study, only heat-induced a0 /21 0 0 APBs and slip-induced a0 /21 1 1 APBs are visible in tensile specimens tested at room temperature. (3) Hot spinning forging, which has the function of refining grains and substructures, combined with subsequent annealing and the selection of suitable alloying elements is a very suitable method to improve the ductility of brittle intermetallic compounds. As a result, the tensile elongation of Fe3 Al at room temperature can be increased to about 10%. Acknowledgement The authors would like to acknowledge Mr. Zheng-Zhong Lei of Division of Research and Development, Tianjin Institute of Metallurgy, for his help on spinning forging performance. References [1] V.K. Sikka, in: L.A. Johnson, D.P. Pope, J.O. Stiegler (Eds.), High Temperature Ordered Intermetallic Alloys IV, Materials Research Society, Pittsburgh, PA, 1991, pp. 907–912. [2] C.G. McKamey, J.H. DeVan, P.F. Tortorelli, V.K. Sikka, J. Mater. Res. 6 (1991) 1779–1805. [3] N.S. Stoloff, Mater. Sci. Eng. A 258 (1998) 1–14. [4] M.G. Mendiratta, S.K. Ehlers, D.K. Chatterjee, H.A. Lipsitt, Metall. Trans. A18 (1987) 283–291. [5] C.G. McKamey, J.A. Horton, C.T. Liu, Scripta Metall. 22 (1988) 1679–1681.
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