- Email: [email protected]
Microstructure and mechanical properties of high Nb containing TiAl alloy parts fabricated by metal injection molding
Materials Science and Engineering A 526 (2009) 31–37
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructure and mechanical properties of high Nb containing TiAl alloy parts fabricated by metal injection molding Haoming Zhang a,∗ , Xinbo He a , Xuanhui Qu a,b , Liming Zhao a a b
School of Materials Science and Engineering, University of Science and Technology Beijing, Beijing, 100083, PR China State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing 100083, PR China
a r t i c l e
i n f o
Article history: Received 11 May 2009 Received in revised form 23 June 2009 Accepted 2 July 2009
Keywords: TiAl based alloy Metal injection molding Microstructure Mechanical property
a b s t r a c t Metal injection molding (MIM) process was applied to fabricate parts of high Nb containing TiAl alloys with a nominal composition of Ti–45Al–8.5Nb–0.2W–0.2B–0.02Y (at.%), and the effects of sintering parameters on their microstructures and mechanical properties, as well as the fractographies after tensile tests were investigated. Results show that for sintering of the alloy in vacuum, effective densification took place in the temperature range 1460–1480 ◦ C. Sintering at too high a temperature or too long a time will result in distortion or warpage of the sintered body or coarsening of the lamellar colony. When the optimum sintering parameters (1480 ◦ C, 2 h) were chosen, the alloy with the relative density of 96.2% was obtained. The microstructure was homogenous and fine-grained near lamellar structure, consisting of ␣2 /␥ lamellar colonies with an average size of 60 m, small amounts of  phase, few boride rods and yttrium oxide precipitates. Its compressive strength, compressibility, ultimate tensile strength and plastic elongation were 2839 MPa, 34.9%, 382 MPa and 0.46%, respectively. At tensile tests, translamellar fracture was the predominant mode and the microcracks often originated from pores and the interfaces of borides/matrix and  phase/matrix. © 2009 Elsevier B.V. All rights reserved.
1. Introduction TiAl based alloys have attracted much attention and have been applied to structural and engine parts of airplanes as well as to automobile engine parts because they have low density, excellent high temperature strength and oxidation resistance [1–3]. Of late, it has been found that Nb is the essential and effective element to improve mechanical properties, especially, high temperature strength for TiAl alloys and the high Nb containing TiAl alloys are considered as the potential high temperature structural materials [4,5]. However, they suffer from poor ductility and toughness at ambient temperature, and the addition of high Nb also increases the difficulties of their fabrication [6,7]. Generally, these alloys are fabricated by ingot metallurgy [8,9], but the poor room temperature ductility causes the subsequently conventional manufacturing operations to be difficult for preparing of high Nb containing TiAl alloy parts. Powder metallurgy (PM) technique appears to be more attractive since high degrees of chemical homogeneities can be obtained and macrosegregations are avoided. Metal injection molding (MIM) is a net-shape powder metallurgy forming process, and is cost-effective for producing small, complex, precision parts in high
∗ Corresponding author. Tel.: +86 10 82377296; fax: +86 10 62334311. E-mail address: [email protected] (H. Zhang). 0921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2009.07.003
volume [10–12]. In this study, the MIM process was applied to the fabrication of high Nb containing TiAl alloy parts, and the suitable injection molding process variables, especially sintering parameters, were obtained. The effects of sintering temperature, time on their microstructures and mechanical properties were investigated. Additionally, their fractographies after tensile tests were also studied. 2. Experimental 2.1. Raw materials High Nb containing TiAl alloy powder with the nominal composition of Ti–45Al–8.5Nb–0.2W–0.2B–0.02Y (at.%) (Ti–45Al–8.5Nb–(W, B, Y)) was used. The prior ingot was fabricated using plasma arc furnace twice and annealed at 1200 ◦ C for 50 h in order to reduce the composition inhomogeneity. The homogeneous ingot was machined to the size of Ø18 mm × 220 mm and then used to produce alloy powder by argon gas atomization process using the PIGA-technique (plasma melting induction guiding gas atomization) at Shanxi Bangzhen Co., Ltd. (China). Details of this technique are given in Ref. [13]. The alloy powders were classified by sieves, and the finer powder with −200 mesh was extracted for the injection molding. Its morphology was shown in Fig. 1. The powder particles with the mean size approxi-
32
H. Zhang et al. / Materials Science and Engineering A 526 (2009) 31–37
Fig. 1. The morphology of Ti–45Al–8.5Nb–(W, B, Y) alloy powder (−200 mesh).
mately 50 m are predominantly spherical and satellite formation is hardly observed, indicating that they are available for the MIM. 2.2. MIM process Details of the fabrication and the parameters in MIM process were given as follows: Feedstock preparation: The feedstock is a mixture of alloy powder and binder at a given ratio. For Ti–45Al–8.5Nb–(W, B, Y) alloy powder, the binder of 63% (wt.%) PW (paraffin wax), 20% (wt.%) LDPE (low density polyethylene), 12% (wt.%) PP (polypropylene) and 5% (wt.%) SA (stearic acid) was found to be suitable. Based on the characteristics and rheological properties of the powder and binder, a feedstock with powder loading of 65% was chosen, which was
Fig. 2. Relationship between relative density and sintering temperature.
prepared by mixing the powder and binder in a SK-160 mixer at a temperature of 140 ◦ C for 1.5 h. Injection molding: Injection compacts were fabricated in a CJ-ZZ injection molding machine. Compacts of complex shaped as well as cuboids with dimensions 35 mm × 20 mm × 6 mm for tests were molded. An injection pressure of 90 MPa, a injection temperature of 150 ◦ C and a mould temperature of 30 ◦ C were turned out to be better parameters for the injection process. Debinding: A two-step debinding of solvent debinding and thermal debinding is necessary for the selected binder system. The solvent debinding was performed in a bath of solvent with trichloroethylene at 40 ◦ C for 12 h. The compacts were taken out and dried for 6 h after bathing. Then thermal debinding was performed in a temperature range between 30 and 600 ◦ C under vacuum atmosphere (10−3 Pa).
Fig. 3. Microstructures of the specimens sintered at different temperatures for 1 h: (a) 1400 ◦ C; (b) 1430 ◦ C; (c) 1460 ◦ C; (d) 1480 ◦ C.
H. Zhang et al. / Materials Science and Engineering A 526 (2009) 31–37
33
3. Results and discussion 3.1. Effect of sintering temperature on microstructure
Fig. 4. XRD patters of the specimens sintered at different temperatures for 1 h: (a) 1400 ◦ C; (b) 1430 ◦ C; (c) 1480 ◦ C.
Table 1 EDS analysis in Fig. 3d. Analysis location
Ti (at.%)
Al (at.%)
Nb (at.%)
W (at.%)
Lamellar colony  phase
46.49 53.72
45.36 34.73
8.12 11.02
− 0.48
Sintering: Sintering was performed in a vacuum atmosphere (10−3 Pa) furnace with molybdenum heating elements. The debinded compacts were sintered at different temperatures ranging from 1400 to 1500 ◦ C and hold for different times ranging from 0.5 to 3 h respectively. 2.3. Characterization The densities of the specimens were measured by the underwater gravimetric method. Microstructural observation was carried out by scanning electron microscopy using back scattering electron (BSE) imaging and energy dispersive spectroscopy (EDS). Constituent phases were characterized by the X-ray diffraction (XRD), which was conducted on a conventional diffractometer using Cu K␣ radiation. Tensile and compressive tests were conducted at room temperature on an Instron material tester at the strain rate of 1 × 10−3 and 5 × 10−3 s−1 , respectively. The specimens were prepared by spark erosion in the form of plate with gauge size of 25 mm × 3 mm × 1.5 mm for tensile tests and of column with gauge size of Ø3 mm × 7 mm for compressive tests. In the gauge region, the specimens were additionally ground with emery paper.
The relative densities of the specimens sintered in vacuum for 1 h at different temperatures were shown in Fig. 2. It was observed that the densities of the specimens increased with the sintering temperature, and the effective densification took place in the temperature range 1460–1480 ◦ C. The maximal relative density of 96% was achieved in the specimen sintered at 1480 ◦ C. With the further increase of the sintering temperature, the densities of the specimens kept little change. Fig. 3 shows the microstructures of the specimens formed at different sintering temperatures for 1 h. Three phases, ␣2 , ␥ and , in the specimens can be identified by X-ray diffraction spectra shown in Fig. 4. Alloy powders are deformed and bonded, and the interfaces of these particles are very evident in the sintered specimen at 1400 ◦ C, as can be seen in Fig. 3a. It can also be found that a large amount of pores and  phase (white phase) exist at prior particle boundaries. When sintering temperature is up to 1430 ◦ C, the interfaces of these particles are undiscernible, and small amounts of  phase are observed at triple junctions of lamellar colonies (Fig. 3b). The microstructure of the specimen sintered at 1460 ◦ C (Fig. 3c) is similar to that sintered at 1480 ◦ C (Fig. 3d) except for the latter having fewer pores and the average size of the lamellar colonies increasing from 40 to 60 m approximately. As shown in Fig. 3d, the homogenous and fine-grained near lamellar microstructure is successfully formed at 1480 ◦ C, which consists of ␣2 /␥ lamellar colonies and  phase. The  phase exists at the triple junctions or the boundaries between the lamellar colonies and many  phases in the forms of stripes are jointed to ring appearing the reticulate morphology. The results of EDS analysis in Table 1 indicate that the  phase contains higher Ti, Nb and W than the nominal composition of the alloy. In addition, there are a few boride particles in the forms of rods and yttrium oxide precipitates in the near lamellar microstructure, and their morphologies in SEM-BSE image from the etched specimens were shown in Fig. 5. Details about their formation mechanism and morphologies in TiAl based alloys can be referred in the literature [14–17]. Fig. 6 shows some pictures of the alloy parts with complex shapes fabricated by MIM process. It must be pointed out that when the sintering temperature was higher than 1480 ◦ C, the geometric shape of the sintered bodies often exhibited gross distortion or warpage, as can be seen in Fig. 6b. Due to the composition of the alloy is not homogeneous completely, when the sintering temperature is too higher, some parts of the sintered body can be melted, and if the liquid phase is too many, it will lead to the distortion of the
Fig. 5. SEM morphologies of borides and yttrium oxides in the etched specimens: (a) rod borides; (b) particle yttrium oxides and rod borides (section imagine).
34
H. Zhang et al. / Materials Science and Engineering A 526 (2009) 31–37
Fig. 6. Pictures of Ti–45Al–8.5Nb–(W, B, Y) alloy parts with complex shapes fabricated by MIM process: (a) sintered at 1480 ◦ C; (b) sintered at 1490 ◦ C (above three) or 1500 ◦ C (below three). Fig. 7. Relationship between relative density and sintering time.
sintered body. Therefore, it can be deduced that 1480 ◦ C is an appropriate sintering temperature to fabricate Ti–45Al–8.5Nb–(W, B, Y) alloys with near lamellar microstructure consisting of homogenous and fine-grained ␣2 /␥ lamellar colonies. Compared with as-cast ingot, which shows microstructure with obvious segregation, such as L-segregation (the scale of the region is about 140–300 m, the distance between L-segregations regions is about 500 m) [18], the MIM alloy showed more homogeneous microstructure and finer average size of colony without macrosegregations, which are beneficial to the improvement of the properties. 3.2. Effect of sintering time on microstructure Fig. 7 shows the relative densities of the specimens sintered in vacuum at 1480 ◦ C for different times. The relative density of the sintered specimen at 1480 ◦ C for 0.5 h was a lower value of 93.1%,
Table 2 EDS analysis in Fig. 8d. Analysis location
Ti (at.%)
Al (at.%)
Nb (at.%)
Lamellar colony Black region
46.35 44.89
45.56 48.34
8.09 6.76
then it rapidly increased to 96% at 1480 ◦ C for 1 h. When sintering time was over 1 h, the relative densities of the specimens only had a slightly increase, i.e. 96.2% for 2 h and 96.3% for 3 h, so the full density cannot be achieved with resort to increased sintering time. Microstructures of the specimens sintered at 1480 ◦ C for different times were shown in Fig. 8. It can be seen that the MIM alloys have near lamellar microstructures for all the different sintering times, but the morphology of the  phase and the lamellar colony
Fig. 8. Microstructures of the specimens sintered at 1480 ◦ C for different times: (a) 0.5 h; (b) 1 h; (c) 2 h; (d) 3 h.
H. Zhang et al. / Materials Science and Engineering A 526 (2009) 31–37 Table 3 Results of compressive and tensile tests at room temperature conducted on the specimens sintered at different parameters. Sintering parameters
Mechanical properties bc (MPa)
1400 ◦ C 1 h 1430 ◦ C 1 h 1460 ◦ C 1 h 1480 ◦ C 0.5 h 1480 ◦ C 1 h 1480 ◦ C 2 h 1480 ◦ C 3 h
1449 2078 2376 2576 2754 2839 2697
± ± ± ± ± ± ±
12 7 8 3 4 4 6
εc (%) 18.2 27.2 32.6 33.5 34.3 34.9 34.1
± ± ± ± ± ± ±
2.4 1.8 1.4 1.6 0.9 0.8 1.1
UTS (MPa)
ı (%)
± ± ± ± ± ± ±
0.19 0.28 0.40 0.42 0.43 0.46 0.38
182 248 312 324 343 382 336
10 7 3 2 2 3 4
± ± ± ± ± ± ±
0.07 0.11 0.09 0.07 0.04 0.03 0.12
bc : ultimate compressive strength; εc : compressibility; UTS: ultimate tensile strength; ı: plastic elongation.
size are different. As shown in Fig. 8a, the microstructure of the specimen sintered at 1480 ◦ C for 0.5 h is similar to that sintered at 1480 ◦ C for 1 h (Fig. 8b), but the amount of the pores are more than the latter. The average size of lamellar colonies is approximately 50 m. When sintering time is extended to 2 h, the  phase
35
appearing reticulate morphology becomes lighter and wider, as can be seen in Fig. 8c, which indicates that the  phase partly transforms to ␣2 and ␥, and the composition of the sintered body becomes more homogenous. In addition, the lamellar colonies size hardly have any increase compared with the specimen sintered at 1480 ◦ C for 1 h. Fig. 8d shows the microstructure of the specimen for sintering 3 h, it was found that the  phase with reticulate morphology cannot be observed, but it changes into the slightly white nubbly segregation distributed in the lamellar colonies. It was also found that the black regions with curve-shaped appear in the boundaries of lamellar colonies. The results of EDS analysis in Table 2 reveal that the black region contains higher Al (3% or so) and lower Nb than the area of lamellar colony, which is similar to the S-segregation described by other investigator in as-cast alloys [19]. It must be noting that the lamellar colonies are severely coarsening in the sintered specimen for 3 h, with the average size of 200 m or so. Additionally, it was also clearly seen that the boride rods and yttrium oxide particles in the sintered bodies cannot be eliminated via extending the sintering time. Hence, for our studied alloy, it seems reasonable that the optimum sintering parameters were chosen as 1480 ◦ C, 2 h.
Fig. 9. SEM images of the fracture surfaces after tensile tests conducted on the specimens sintered at different temperatures: (a, b) 1400 ◦ C; (c, d) 1460 ◦ C; (e, f) 1480 ◦ C.
36
H. Zhang et al. / Materials Science and Engineering A 526 (2009) 31–37
Fig. 10. SEM images of segregations on the tensile fracture surface: (a and b) boride and its indentation; (c and d)  phase and its indentation.
3.3. Mechanical properties The results of compressive and tensile tests at room temperature conducted on the specimens sintered at different parameters were listed in Table 3. It can be clearly seen that the specimen sintered at 1480 ◦ C for 2 h has the best properties than the others. This is in accordance with the microstructure of it, which has fewer voids or pores than other specimens and the finer grain size. In general, the voids or pores are often acted as microcracks under tensile or compressive loading, and a material with fine microstructure possesses high yield strength according to Hall–Petch’s Law. It was also found that there exhibits huge differences between the tensile and compressive properties in our studied alloy. From the data, the bc is 6–8 times UTS, and the εc is 50–60 times ı. This is correlated with the brittle characteristic inherent in TiAl based alloys. Under tensile loading, the microcracks generate at a lower loading and the driving force of the crack extension is the tensile stress, not the shearing stress or anelastic strain. So, if the cracks come forth, it will rapidly propagate leading to the specimens splitting immediately. Whereas under compressive loading, even though the cracks come forth, it will not be propagating too rapidly. It must be pointed out that although the microstructures of MIM alloys exhibit the finer lamellar colony size and without macrosegregations compared with as-cast alloys, the values of the UTS and ı at room temperature are still below as-cast alloys reported up to date [20]. The reasons in all probability attribute to the existence of higher porosity level and higher oxygen content in the MIM alloys. As an example, the specimen sintered 1480 ◦ C for 2 h exhibits nearly 4% porosity and about 1800 ppm oxygen content, which are higher than that of as-cast alloys [9]. The high oxygen content in MIM alloys are not difficult to understand when the MIM process is taken into consideration. In MIM process, the raw alloying powders with higher mobilities can be severely contaminated by oxygen and the binder system during feedstock preparation, debinding and sinter-
ing. The extremely high interstitial oxygen content in ␣2 (Ti3 Al) and ␥ (TiAl) phases can severely embrittle the MIM alloys. Therefore, it is believed that the MIM alloys can exhibit excellent mechanical properties if the oxygen content and flaw density are reduced. 3.4. Fractographies after tensile tests Fig. 9 shows the morphologies of tensile fracture surfaces of the specimens sintered at different temperatures. For the specimen at 1400 ◦ C, there are many, irregular and connected pores in their fracture surface (Fig. 9a and b). Its fracture mode shows brittle fracture, and the structure nearby the pores can easily be fractured during tensile loading, leading to the lower strength of the specimen sintered at 1400 ◦ C. With the sintering time increasing, the pores in the specimens decrease markedly. For the specimens at 1460 and 1480 ◦ C, the fracture mode is brittle exhibiting translamellar fracture predominantly (Fig. 9c and e), and in some areas, there also exhibit interlamellar fracture, as can be seen in Fig. 9d and f. It is worth noting that the boride rods and strip-shaped  phase segregations are frequently observed on the fracture surfaces, as shown by A and B arrows in Fig. 10a and c, respectively. The EDS analysis was performed to identify them, and its results were listed in Table 4. As shown by arrow in Fig. 10b, the boride rod has been pulled out of the metal matrix during failure, leaving the indentation on the fracture surface, and the inside surface of the indentation is very smooth, indicating a debonding process along the interface between boride rod and the matrix. From Fig. 10d, it also can be seen that the strip-shaped indentation is in the crossing area of the Table 4 EDS analysis in Fig. 10a and c. Analysis location
Ti (at.%)
Al (at.%)
Nb (at.%)
W (at.%)
B (at.%)
A (boride) B ( phase)
21.43 54.71
12.12 33.26
3.78 11.64
− 0.37
62.55 −
H. Zhang et al. / Materials Science and Engineering A 526 (2009) 31–37
two surfaces, where the prior  phase often exists. For our studied alloy, although more rigorous investigations and studies are needed for the effects of boride and  phase segregations on the mechanical properties, it can be preliminarily concluded that that the deformation incompatibilities of the boride/matrix and the  phase/matrix will cause stress concentration and induce microcracks between their interfaces, which have negative effects to the mechanical properties. 4. Conclusions The complex-shaped parts of Ti–45Al–8.5Nb–(W, B, Y) alloys were successfully fabricated by MIM process. From the results of the present investigation, the following conclusions can be drawn: (1) For sintering of Ti–45Al–8.5Nb–(W, B, Y) alloy in vacuum, effective deification took place in the temperature range 1460–1480 ◦ C. A full density alloy cannot be obtained via increasing the sintering temperature and time. Sintering at too high a temperature will result in distortion or warpage of the sintered body, and sintering for too long a time will result in coarsening of the lamellar colony. The optimum sintering parameters were chosen as 1480 ◦ C, 2 h. (2) When the optimum sintering parameters were chosen, the alloy with the relative density of 96.2% was obtained. The microstructure was homogenous and fine-grained near lamellar structure, consisting of ␣2 /␥ lamellar colonies with an average size of 60 m,  phase, few boride rods and yttrium oxide precipitates. Its compressive strength, compressibility, ultimate tensile strength and plastic elongation were 2839 MPa, 34.9%, 382 MPa and 0.46%, respectively. (3) The UTS and ı values at room temperature are below as-cast alloys reported up to date. The reason is due to the existence
37
of higher porosity level and higher oxygen content in the MIM alloys. At tensile test, translamellar fracture was the predominant mode and the microcracks often originated from pores and the interfaces of boride/matrix and  phase/matrix. Acknowledgements This research was sponsored by the Key Grant Project of Chinese Ministry of Education (No. 704008) and by the program from New Century Excellent Talents in University (NCET-06-0081). References [1] Y.W. Kim, Acta Metall. 40 (1992) 1121–1134. [2] H. Clemens, Adv. Eng. Mater. 9 (2000) 551–570. [3] Z. Jin, C. Cady, G.T. Gray, Y.W. Kim, Metall. Mater. Trans. A 31 (2000) 1007– 1016. [4] G.L. Chen, L.C. Zhang, Mater. Sci. Eng. A329/331 (2002) 163–170. [5] W.J. Zhang, S.C. Deevi, G.L. Chen, Intermetallics 10 (2002) 403–406. [6] S. Bystrzanowski, A. Bartels, H. Clemens, R. Gerling, F.P. Schimansky, H. Kestler, G. Dehm, G. Haneczok, M. Weller, Intermetallics 13 (2005) 515–524. [7] W.J. Zhang, F. Appel, Mater. Sci. Eng. A329/331 (2002) 649–652. [8] Z.C. Liu, J.P. Lin, S.J. Li, G.L. Chen, Intermetallics 10 (2002) 653–659. [9] X.J. Xu, L.H. Xu, J.P. Lin, Y.L. Wang, Z. Lin, G.L. Chen, Intermetallics 13 (2005) 337–341. [10] J.R. Merhar, Met. Powder Rep. 45 (1990) 339–342. [11] R.M. German, Int. J. Powder Metall. 25 (1993) 165–169. [12] K.M. Kulkarni, Int. J. Powder Metall. 36 (2000) 43–52. [13] R. Gerling, H. Clemens, F.P. Schimansky, Adv. Eng. Mater. 6 (2004) 23–28. [14] D. Hu, J.F. Mei, M. Wickins, R.A. Harding, Scr. Mater. 47 (2002) 273–278. [15] W.J. Zhang, S.C. Deevi, Mater. Sci. Eng. A 337 (2002) 17–20. [16] D. Hu, Intermetallics 10 (2002) 851–858. [17] Y.Q. Yan, L.Q. Zhang, G.Z. Lou, K.G. Wang, L. Zhou, Mater. Sci. Eng. A 280 (2000) 187–191. [18] X.J. Xu, J.P. Lin, Z.K. Teng, Y.L. Wang, G.L. Chen, Mater. Lett. 61 (2007) 369–373. [19] G.L. Chen, X.J. Xu, Z.K. Teng, Y.L. Wang, J.P. Lin, Intermetallics 15 (2007) 625–631. [20] X.J. Xu, J.P. Lin, Y.L. Wang, J.F. Gao, Z. Lin, G.L. Chen, J. Alloys Compd. 414 (2006) 131–136.
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructure and mechanical properties of high Nb containing TiAl alloy parts fabricated by metal injection molding Haoming Zhang a,∗ , Xinbo He a , Xuanhui Qu a,b , Liming Zhao a a b
School of Materials Science and Engineering, University of Science and Technology Beijing, Beijing, 100083, PR China State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing 100083, PR China
a r t i c l e
i n f o
Article history: Received 11 May 2009 Received in revised form 23 June 2009 Accepted 2 July 2009
Keywords: TiAl based alloy Metal injection molding Microstructure Mechanical property
a b s t r a c t Metal injection molding (MIM) process was applied to fabricate parts of high Nb containing TiAl alloys with a nominal composition of Ti–45Al–8.5Nb–0.2W–0.2B–0.02Y (at.%), and the effects of sintering parameters on their microstructures and mechanical properties, as well as the fractographies after tensile tests were investigated. Results show that for sintering of the alloy in vacuum, effective densification took place in the temperature range 1460–1480 ◦ C. Sintering at too high a temperature or too long a time will result in distortion or warpage of the sintered body or coarsening of the lamellar colony. When the optimum sintering parameters (1480 ◦ C, 2 h) were chosen, the alloy with the relative density of 96.2% was obtained. The microstructure was homogenous and fine-grained near lamellar structure, consisting of ␣2 /␥ lamellar colonies with an average size of 60 m, small amounts of  phase, few boride rods and yttrium oxide precipitates. Its compressive strength, compressibility, ultimate tensile strength and plastic elongation were 2839 MPa, 34.9%, 382 MPa and 0.46%, respectively. At tensile tests, translamellar fracture was the predominant mode and the microcracks often originated from pores and the interfaces of borides/matrix and  phase/matrix. © 2009 Elsevier B.V. All rights reserved.
1. Introduction TiAl based alloys have attracted much attention and have been applied to structural and engine parts of airplanes as well as to automobile engine parts because they have low density, excellent high temperature strength and oxidation resistance [1–3]. Of late, it has been found that Nb is the essential and effective element to improve mechanical properties, especially, high temperature strength for TiAl alloys and the high Nb containing TiAl alloys are considered as the potential high temperature structural materials [4,5]. However, they suffer from poor ductility and toughness at ambient temperature, and the addition of high Nb also increases the difficulties of their fabrication [6,7]. Generally, these alloys are fabricated by ingot metallurgy [8,9], but the poor room temperature ductility causes the subsequently conventional manufacturing operations to be difficult for preparing of high Nb containing TiAl alloy parts. Powder metallurgy (PM) technique appears to be more attractive since high degrees of chemical homogeneities can be obtained and macrosegregations are avoided. Metal injection molding (MIM) is a net-shape powder metallurgy forming process, and is cost-effective for producing small, complex, precision parts in high
∗ Corresponding author. Tel.: +86 10 82377296; fax: +86 10 62334311. E-mail address: [email protected] (H. Zhang). 0921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2009.07.003
volume [10–12]. In this study, the MIM process was applied to the fabrication of high Nb containing TiAl alloy parts, and the suitable injection molding process variables, especially sintering parameters, were obtained. The effects of sintering temperature, time on their microstructures and mechanical properties were investigated. Additionally, their fractographies after tensile tests were also studied. 2. Experimental 2.1. Raw materials High Nb containing TiAl alloy powder with the nominal composition of Ti–45Al–8.5Nb–0.2W–0.2B–0.02Y (at.%) (Ti–45Al–8.5Nb–(W, B, Y)) was used. The prior ingot was fabricated using plasma arc furnace twice and annealed at 1200 ◦ C for 50 h in order to reduce the composition inhomogeneity. The homogeneous ingot was machined to the size of Ø18 mm × 220 mm and then used to produce alloy powder by argon gas atomization process using the PIGA-technique (plasma melting induction guiding gas atomization) at Shanxi Bangzhen Co., Ltd. (China). Details of this technique are given in Ref. [13]. The alloy powders were classified by sieves, and the finer powder with −200 mesh was extracted for the injection molding. Its morphology was shown in Fig. 1. The powder particles with the mean size approxi-
32
H. Zhang et al. / Materials Science and Engineering A 526 (2009) 31–37
Fig. 1. The morphology of Ti–45Al–8.5Nb–(W, B, Y) alloy powder (−200 mesh).
mately 50 m are predominantly spherical and satellite formation is hardly observed, indicating that they are available for the MIM. 2.2. MIM process Details of the fabrication and the parameters in MIM process were given as follows: Feedstock preparation: The feedstock is a mixture of alloy powder and binder at a given ratio. For Ti–45Al–8.5Nb–(W, B, Y) alloy powder, the binder of 63% (wt.%) PW (paraffin wax), 20% (wt.%) LDPE (low density polyethylene), 12% (wt.%) PP (polypropylene) and 5% (wt.%) SA (stearic acid) was found to be suitable. Based on the characteristics and rheological properties of the powder and binder, a feedstock with powder loading of 65% was chosen, which was
Fig. 2. Relationship between relative density and sintering temperature.
prepared by mixing the powder and binder in a SK-160 mixer at a temperature of 140 ◦ C for 1.5 h. Injection molding: Injection compacts were fabricated in a CJ-ZZ injection molding machine. Compacts of complex shaped as well as cuboids with dimensions 35 mm × 20 mm × 6 mm for tests were molded. An injection pressure of 90 MPa, a injection temperature of 150 ◦ C and a mould temperature of 30 ◦ C were turned out to be better parameters for the injection process. Debinding: A two-step debinding of solvent debinding and thermal debinding is necessary for the selected binder system. The solvent debinding was performed in a bath of solvent with trichloroethylene at 40 ◦ C for 12 h. The compacts were taken out and dried for 6 h after bathing. Then thermal debinding was performed in a temperature range between 30 and 600 ◦ C under vacuum atmosphere (10−3 Pa).
Fig. 3. Microstructures of the specimens sintered at different temperatures for 1 h: (a) 1400 ◦ C; (b) 1430 ◦ C; (c) 1460 ◦ C; (d) 1480 ◦ C.
H. Zhang et al. / Materials Science and Engineering A 526 (2009) 31–37
33
3. Results and discussion 3.1. Effect of sintering temperature on microstructure
Fig. 4. XRD patters of the specimens sintered at different temperatures for 1 h: (a) 1400 ◦ C; (b) 1430 ◦ C; (c) 1480 ◦ C.
Table 1 EDS analysis in Fig. 3d. Analysis location
Ti (at.%)
Al (at.%)
Nb (at.%)
W (at.%)
Lamellar colony  phase
46.49 53.72
45.36 34.73
8.12 11.02
− 0.48
Sintering: Sintering was performed in a vacuum atmosphere (10−3 Pa) furnace with molybdenum heating elements. The debinded compacts were sintered at different temperatures ranging from 1400 to 1500 ◦ C and hold for different times ranging from 0.5 to 3 h respectively. 2.3. Characterization The densities of the specimens were measured by the underwater gravimetric method. Microstructural observation was carried out by scanning electron microscopy using back scattering electron (BSE) imaging and energy dispersive spectroscopy (EDS). Constituent phases were characterized by the X-ray diffraction (XRD), which was conducted on a conventional diffractometer using Cu K␣ radiation. Tensile and compressive tests were conducted at room temperature on an Instron material tester at the strain rate of 1 × 10−3 and 5 × 10−3 s−1 , respectively. The specimens were prepared by spark erosion in the form of plate with gauge size of 25 mm × 3 mm × 1.5 mm for tensile tests and of column with gauge size of Ø3 mm × 7 mm for compressive tests. In the gauge region, the specimens were additionally ground with emery paper.
The relative densities of the specimens sintered in vacuum for 1 h at different temperatures were shown in Fig. 2. It was observed that the densities of the specimens increased with the sintering temperature, and the effective densification took place in the temperature range 1460–1480 ◦ C. The maximal relative density of 96% was achieved in the specimen sintered at 1480 ◦ C. With the further increase of the sintering temperature, the densities of the specimens kept little change. Fig. 3 shows the microstructures of the specimens formed at different sintering temperatures for 1 h. Three phases, ␣2 , ␥ and , in the specimens can be identified by X-ray diffraction spectra shown in Fig. 4. Alloy powders are deformed and bonded, and the interfaces of these particles are very evident in the sintered specimen at 1400 ◦ C, as can be seen in Fig. 3a. It can also be found that a large amount of pores and  phase (white phase) exist at prior particle boundaries. When sintering temperature is up to 1430 ◦ C, the interfaces of these particles are undiscernible, and small amounts of  phase are observed at triple junctions of lamellar colonies (Fig. 3b). The microstructure of the specimen sintered at 1460 ◦ C (Fig. 3c) is similar to that sintered at 1480 ◦ C (Fig. 3d) except for the latter having fewer pores and the average size of the lamellar colonies increasing from 40 to 60 m approximately. As shown in Fig. 3d, the homogenous and fine-grained near lamellar microstructure is successfully formed at 1480 ◦ C, which consists of ␣2 /␥ lamellar colonies and  phase. The  phase exists at the triple junctions or the boundaries between the lamellar colonies and many  phases in the forms of stripes are jointed to ring appearing the reticulate morphology. The results of EDS analysis in Table 1 indicate that the  phase contains higher Ti, Nb and W than the nominal composition of the alloy. In addition, there are a few boride particles in the forms of rods and yttrium oxide precipitates in the near lamellar microstructure, and their morphologies in SEM-BSE image from the etched specimens were shown in Fig. 5. Details about their formation mechanism and morphologies in TiAl based alloys can be referred in the literature [14–17]. Fig. 6 shows some pictures of the alloy parts with complex shapes fabricated by MIM process. It must be pointed out that when the sintering temperature was higher than 1480 ◦ C, the geometric shape of the sintered bodies often exhibited gross distortion or warpage, as can be seen in Fig. 6b. Due to the composition of the alloy is not homogeneous completely, when the sintering temperature is too higher, some parts of the sintered body can be melted, and if the liquid phase is too many, it will lead to the distortion of the
Fig. 5. SEM morphologies of borides and yttrium oxides in the etched specimens: (a) rod borides; (b) particle yttrium oxides and rod borides (section imagine).
34
H. Zhang et al. / Materials Science and Engineering A 526 (2009) 31–37
Fig. 6. Pictures of Ti–45Al–8.5Nb–(W, B, Y) alloy parts with complex shapes fabricated by MIM process: (a) sintered at 1480 ◦ C; (b) sintered at 1490 ◦ C (above three) or 1500 ◦ C (below three). Fig. 7. Relationship between relative density and sintering time.
sintered body. Therefore, it can be deduced that 1480 ◦ C is an appropriate sintering temperature to fabricate Ti–45Al–8.5Nb–(W, B, Y) alloys with near lamellar microstructure consisting of homogenous and fine-grained ␣2 /␥ lamellar colonies. Compared with as-cast ingot, which shows microstructure with obvious segregation, such as L-segregation (the scale of the region is about 140–300 m, the distance between L-segregations regions is about 500 m) [18], the MIM alloy showed more homogeneous microstructure and finer average size of colony without macrosegregations, which are beneficial to the improvement of the properties. 3.2. Effect of sintering time on microstructure Fig. 7 shows the relative densities of the specimens sintered in vacuum at 1480 ◦ C for different times. The relative density of the sintered specimen at 1480 ◦ C for 0.5 h was a lower value of 93.1%,
Table 2 EDS analysis in Fig. 8d. Analysis location
Ti (at.%)
Al (at.%)
Nb (at.%)
Lamellar colony Black region
46.35 44.89
45.56 48.34
8.09 6.76
then it rapidly increased to 96% at 1480 ◦ C for 1 h. When sintering time was over 1 h, the relative densities of the specimens only had a slightly increase, i.e. 96.2% for 2 h and 96.3% for 3 h, so the full density cannot be achieved with resort to increased sintering time. Microstructures of the specimens sintered at 1480 ◦ C for different times were shown in Fig. 8. It can be seen that the MIM alloys have near lamellar microstructures for all the different sintering times, but the morphology of the  phase and the lamellar colony
Fig. 8. Microstructures of the specimens sintered at 1480 ◦ C for different times: (a) 0.5 h; (b) 1 h; (c) 2 h; (d) 3 h.
H. Zhang et al. / Materials Science and Engineering A 526 (2009) 31–37 Table 3 Results of compressive and tensile tests at room temperature conducted on the specimens sintered at different parameters. Sintering parameters
Mechanical properties bc (MPa)
1400 ◦ C 1 h 1430 ◦ C 1 h 1460 ◦ C 1 h 1480 ◦ C 0.5 h 1480 ◦ C 1 h 1480 ◦ C 2 h 1480 ◦ C 3 h
1449 2078 2376 2576 2754 2839 2697
± ± ± ± ± ± ±
12 7 8 3 4 4 6
εc (%) 18.2 27.2 32.6 33.5 34.3 34.9 34.1
± ± ± ± ± ± ±
2.4 1.8 1.4 1.6 0.9 0.8 1.1
UTS (MPa)
ı (%)
± ± ± ± ± ± ±
0.19 0.28 0.40 0.42 0.43 0.46 0.38
182 248 312 324 343 382 336
10 7 3 2 2 3 4
± ± ± ± ± ± ±
0.07 0.11 0.09 0.07 0.04 0.03 0.12
bc : ultimate compressive strength; εc : compressibility; UTS: ultimate tensile strength; ı: plastic elongation.
size are different. As shown in Fig. 8a, the microstructure of the specimen sintered at 1480 ◦ C for 0.5 h is similar to that sintered at 1480 ◦ C for 1 h (Fig. 8b), but the amount of the pores are more than the latter. The average size of lamellar colonies is approximately 50 m. When sintering time is extended to 2 h, the  phase
35
appearing reticulate morphology becomes lighter and wider, as can be seen in Fig. 8c, which indicates that the  phase partly transforms to ␣2 and ␥, and the composition of the sintered body becomes more homogenous. In addition, the lamellar colonies size hardly have any increase compared with the specimen sintered at 1480 ◦ C for 1 h. Fig. 8d shows the microstructure of the specimen for sintering 3 h, it was found that the  phase with reticulate morphology cannot be observed, but it changes into the slightly white nubbly segregation distributed in the lamellar colonies. It was also found that the black regions with curve-shaped appear in the boundaries of lamellar colonies. The results of EDS analysis in Table 2 reveal that the black region contains higher Al (3% or so) and lower Nb than the area of lamellar colony, which is similar to the S-segregation described by other investigator in as-cast alloys [19]. It must be noting that the lamellar colonies are severely coarsening in the sintered specimen for 3 h, with the average size of 200 m or so. Additionally, it was also clearly seen that the boride rods and yttrium oxide particles in the sintered bodies cannot be eliminated via extending the sintering time. Hence, for our studied alloy, it seems reasonable that the optimum sintering parameters were chosen as 1480 ◦ C, 2 h.
Fig. 9. SEM images of the fracture surfaces after tensile tests conducted on the specimens sintered at different temperatures: (a, b) 1400 ◦ C; (c, d) 1460 ◦ C; (e, f) 1480 ◦ C.
36
H. Zhang et al. / Materials Science and Engineering A 526 (2009) 31–37
Fig. 10. SEM images of segregations on the tensile fracture surface: (a and b) boride and its indentation; (c and d)  phase and its indentation.
3.3. Mechanical properties The results of compressive and tensile tests at room temperature conducted on the specimens sintered at different parameters were listed in Table 3. It can be clearly seen that the specimen sintered at 1480 ◦ C for 2 h has the best properties than the others. This is in accordance with the microstructure of it, which has fewer voids or pores than other specimens and the finer grain size. In general, the voids or pores are often acted as microcracks under tensile or compressive loading, and a material with fine microstructure possesses high yield strength according to Hall–Petch’s Law. It was also found that there exhibits huge differences between the tensile and compressive properties in our studied alloy. From the data, the bc is 6–8 times UTS, and the εc is 50–60 times ı. This is correlated with the brittle characteristic inherent in TiAl based alloys. Under tensile loading, the microcracks generate at a lower loading and the driving force of the crack extension is the tensile stress, not the shearing stress or anelastic strain. So, if the cracks come forth, it will rapidly propagate leading to the specimens splitting immediately. Whereas under compressive loading, even though the cracks come forth, it will not be propagating too rapidly. It must be pointed out that although the microstructures of MIM alloys exhibit the finer lamellar colony size and without macrosegregations compared with as-cast alloys, the values of the UTS and ı at room temperature are still below as-cast alloys reported up to date [20]. The reasons in all probability attribute to the existence of higher porosity level and higher oxygen content in the MIM alloys. As an example, the specimen sintered 1480 ◦ C for 2 h exhibits nearly 4% porosity and about 1800 ppm oxygen content, which are higher than that of as-cast alloys [9]. The high oxygen content in MIM alloys are not difficult to understand when the MIM process is taken into consideration. In MIM process, the raw alloying powders with higher mobilities can be severely contaminated by oxygen and the binder system during feedstock preparation, debinding and sinter-
ing. The extremely high interstitial oxygen content in ␣2 (Ti3 Al) and ␥ (TiAl) phases can severely embrittle the MIM alloys. Therefore, it is believed that the MIM alloys can exhibit excellent mechanical properties if the oxygen content and flaw density are reduced. 3.4. Fractographies after tensile tests Fig. 9 shows the morphologies of tensile fracture surfaces of the specimens sintered at different temperatures. For the specimen at 1400 ◦ C, there are many, irregular and connected pores in their fracture surface (Fig. 9a and b). Its fracture mode shows brittle fracture, and the structure nearby the pores can easily be fractured during tensile loading, leading to the lower strength of the specimen sintered at 1400 ◦ C. With the sintering time increasing, the pores in the specimens decrease markedly. For the specimens at 1460 and 1480 ◦ C, the fracture mode is brittle exhibiting translamellar fracture predominantly (Fig. 9c and e), and in some areas, there also exhibit interlamellar fracture, as can be seen in Fig. 9d and f. It is worth noting that the boride rods and strip-shaped  phase segregations are frequently observed on the fracture surfaces, as shown by A and B arrows in Fig. 10a and c, respectively. The EDS analysis was performed to identify them, and its results were listed in Table 4. As shown by arrow in Fig. 10b, the boride rod has been pulled out of the metal matrix during failure, leaving the indentation on the fracture surface, and the inside surface of the indentation is very smooth, indicating a debonding process along the interface between boride rod and the matrix. From Fig. 10d, it also can be seen that the strip-shaped indentation is in the crossing area of the Table 4 EDS analysis in Fig. 10a and c. Analysis location
Ti (at.%)
Al (at.%)
Nb (at.%)
W (at.%)
B (at.%)
A (boride) B ( phase)
21.43 54.71
12.12 33.26
3.78 11.64
− 0.37
62.55 −
H. Zhang et al. / Materials Science and Engineering A 526 (2009) 31–37
two surfaces, where the prior  phase often exists. For our studied alloy, although more rigorous investigations and studies are needed for the effects of boride and  phase segregations on the mechanical properties, it can be preliminarily concluded that that the deformation incompatibilities of the boride/matrix and the  phase/matrix will cause stress concentration and induce microcracks between their interfaces, which have negative effects to the mechanical properties. 4. Conclusions The complex-shaped parts of Ti–45Al–8.5Nb–(W, B, Y) alloys were successfully fabricated by MIM process. From the results of the present investigation, the following conclusions can be drawn: (1) For sintering of Ti–45Al–8.5Nb–(W, B, Y) alloy in vacuum, effective deification took place in the temperature range 1460–1480 ◦ C. A full density alloy cannot be obtained via increasing the sintering temperature and time. Sintering at too high a temperature will result in distortion or warpage of the sintered body, and sintering for too long a time will result in coarsening of the lamellar colony. The optimum sintering parameters were chosen as 1480 ◦ C, 2 h. (2) When the optimum sintering parameters were chosen, the alloy with the relative density of 96.2% was obtained. The microstructure was homogenous and fine-grained near lamellar structure, consisting of ␣2 /␥ lamellar colonies with an average size of 60 m,  phase, few boride rods and yttrium oxide precipitates. Its compressive strength, compressibility, ultimate tensile strength and plastic elongation were 2839 MPa, 34.9%, 382 MPa and 0.46%, respectively. (3) The UTS and ı values at room temperature are below as-cast alloys reported up to date. The reason is due to the existence
37
of higher porosity level and higher oxygen content in the MIM alloys. At tensile test, translamellar fracture was the predominant mode and the microcracks often originated from pores and the interfaces of boride/matrix and  phase/matrix. Acknowledgements This research was sponsored by the Key Grant Project of Chinese Ministry of Education (No. 704008) and by the program from New Century Excellent Talents in University (NCET-06-0081). References [1] Y.W. Kim, Acta Metall. 40 (1992) 1121–1134. [2] H. Clemens, Adv. Eng. Mater. 9 (2000) 551–570. [3] Z. Jin, C. Cady, G.T. Gray, Y.W. Kim, Metall. Mater. Trans. A 31 (2000) 1007– 1016. [4] G.L. Chen, L.C. Zhang, Mater. Sci. Eng. A329/331 (2002) 163–170. [5] W.J. Zhang, S.C. Deevi, G.L. Chen, Intermetallics 10 (2002) 403–406. [6] S. Bystrzanowski, A. Bartels, H. Clemens, R. Gerling, F.P. Schimansky, H. Kestler, G. Dehm, G. Haneczok, M. Weller, Intermetallics 13 (2005) 515–524. [7] W.J. Zhang, F. Appel, Mater. Sci. Eng. A329/331 (2002) 649–652. [8] Z.C. Liu, J.P. Lin, S.J. Li, G.L. Chen, Intermetallics 10 (2002) 653–659. [9] X.J. Xu, L.H. Xu, J.P. Lin, Y.L. Wang, Z. Lin, G.L. Chen, Intermetallics 13 (2005) 337–341. [10] J.R. Merhar, Met. Powder Rep. 45 (1990) 339–342. [11] R.M. German, Int. J. Powder Metall. 25 (1993) 165–169. [12] K.M. Kulkarni, Int. J. Powder Metall. 36 (2000) 43–52. [13] R. Gerling, H. Clemens, F.P. Schimansky, Adv. Eng. Mater. 6 (2004) 23–28. [14] D. Hu, J.F. Mei, M. Wickins, R.A. Harding, Scr. Mater. 47 (2002) 273–278. [15] W.J. Zhang, S.C. Deevi, Mater. Sci. Eng. A 337 (2002) 17–20. [16] D. Hu, Intermetallics 10 (2002) 851–858. [17] Y.Q. Yan, L.Q. Zhang, G.Z. Lou, K.G. Wang, L. Zhou, Mater. Sci. Eng. A 280 (2000) 187–191. [18] X.J. Xu, J.P. Lin, Z.K. Teng, Y.L. Wang, G.L. Chen, Mater. Lett. 61 (2007) 369–373. [19] G.L. Chen, X.J. Xu, Z.K. Teng, Y.L. Wang, J.P. Lin, Intermetallics 15 (2007) 625–631. [20] X.J. Xu, J.P. Lin, Y.L. Wang, J.F. Gao, Z. Lin, G.L. Chen, J. Alloys Compd. 414 (2006) 131–136.