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Characterization of Gd-rich precipitates in a fully lamellar TiAl alloy

Characterization of Gd-rich precipitates in a fully lamellar TiAl alloy

Scripta Materialia 137 (2017) 50–54 Contents lists available at ScienceDirect Scripta Materialia journal homepage: www.elsevier.com/locate/scriptama...

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Scripta Materialia 137 (2017) 50–54

Contents lists available at ScienceDirect

Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat

Regular Article

Characterization of Gd-rich precipitates in a fully lamellar TiAl alloy Xiaodong Wang a, Ruichun Luo a, Fang Liu a, Fan Zhu a, Shuangxi Song a, Bin Chen a,⁎, Xiwen Zhang b, Ji Zhang b, Mingwei Chen a,c,d,⁎⁎ a

State Key Laboratory of Metal Matrix Composites, School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200240, China High Temperature Materials Division, Central Iron and Steel Research Institute, Beijing 100081, China Department of Materials Science and Engineering, Johns Hopkins University, Baltimore, MD 21218, USA d Advanced Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan b c

a r t i c l e

i n f o

Article history: Received 22 March 2017 Received in revised form 21 April 2017 Accepted 26 April 2017 Available online xxxx Keywords: TiAl Rare earth element Microstructure refinement Precipitate Characterization

a b s t r a c t We report microstructural characterization of a Gd-modified TiAl alloy with a fully lamellar structure by utilizing Cs-corrected transmission electron microscopy. Two types of Gd-contained precipitates, Gd2O3 and Al2Gd, with sizes ranging from several nanometers to micrometers were detected at grain boundaries and lamellar interfaces. Atomic-scale images reveal a well-defined crystallographic relationship between the nano-scale Al2Gd precipitates and γ-TiAl while Gd2O3, formed from TiAl melt during solidification and precipitated from α phase, is incoherent with TiAl matrix. The microstructural characterization provides direct evidence that the microstructure refinement of the Gd-modified TiAl is mainly from both oxide and intermetallic precipitates. © 2017 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

TiAl alloys have attracted great attention in aerospace and automobile applications due to their low density, high specific strength and stiffness as well as strength retention at elevated temperatures [1,2]. However, the intrinsic poor ductility and lower fracture toughness of TiAl alloys at ambient temperature have limited their applications to a great extent. It is well known that microstructure refinement is one of the most effective ways to improve both strength and ductility of materials. Alloying has been demonstrated to be a viable approach to enhance the room-temperature mechanical properties of TiAl alloys by microstructure refinement [3–5]. However, in comparison with duplex TiAl alloys produced by thermal treatment in the (α + γ) two-phase region, the fully lamellar TiAl alloys, formed by high-temperature treatment in the single α-phase region, have much poor response to alloying addition in microstructure refinement and ductility improvement. The rapid α grain growth at a high temperature above 1300 °C often gives rise to coarsened grains as large as thousands of micrometers and leads to further deterioration of the room-temperature ductility and strength [2,3]. Rare earth elements (REEs) are few additions which remain functional in microstructure refinement of the fully lamellar TiAl [6–11], similar to the role in other alloy systems [12–14]. Previous studies have suggested that the addition of yttrium [7], neodymium [8] and gadolinium [6,9–11] can obviously reduce the grain ⁎ Corresponding author. ⁎⁎ Correspondence to: M. Chen, Advanced Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan. E-mail addresses: [email protected] (B. Chen), [email protected] (M. Chen).

http://dx.doi.org/10.1016/j.scriptamat.2017.04.038 1359-6462/© 2017 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

sizes and lamellar spacings of TiAl alloys and improve both room-temperature strength and ductility as well as high-temperature creep resistance. However, the mechanisms of the microstructure refinement in the REE-modified TiAl alloys have not been well understood because of complicate solidification sequences and solid-state phase transformations of TiAl during thermal treatment and cooling [15–17]. In this study we employed the Cs-corrected transmission electron microscopy (TEM) to systematically investigate the microstructure and chemistry of a Gdmodified TiAl alloy with a fully lamellar structure, unveiling the micromechanisms of the REE effects in the microstructure refinement of fully lamellar TiAl. The alloy used in this study has a nominal compositions of Ti-46.5Al2.5 V-1.0Cr-0.3Ni-0.15Gd (at.%). Additionally, a Gd-free Ti-46.5Al-2.5 V1.0Cr-0.3Ni alloy was also prepared as the reference. Ingots were prepared by cold crucible induction levitation melting, cast into 85 mm diameter graphite permanent mold, and subsequently hot isostatically pressed. The cylindrical ingots were isothermally forged into plates at 1050 °C with 85% height reduction. The samples used in this study were taken from the homogenous zones of the as-forged plates with dimensions approximate 10 × 10 × 15 mm and then sealed in quartz tubes back filled with Ar for a series of heat treatments. The samples were solution treated in the single α phase region at 1340 °C for 0.5 h, 1 h and 2 h, respectively, followed by air cooling to room temperature. A fully lamellar microstructure was formed during the cooling by a eutectoid reaction from α-TiAl to γ-TiAl and α2-Ti3Al. Samples for optical microscope (OM) observations were prepared by following the standard metallography procedure. The polished

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surface was etched in a hydro-solution containing 2 vol% HF and 10 vol% HNO3. The as-polished samples were also inspected by a scanning electron microscope (SEM, JEOL JEM-7800F) to obtain back scattered electron (BSE) images. TEM specimens were prepared by twin-jet electropolishing with a electrolyte containing 10 vol% perchloric acid, 70 vol% butanol and 20 vol% ethanol at −30 °C. TEM and high-angle annular dark-field scanning transmission electron microscopy (HAADFSTEM) observations were conducted using a JEOL ARM-200F equipped with a Cs corrector, operated at an acceleration voltage of 200 kV. Energy Dispersive Spectrometer (EDS) elemental mappings were acquired using a JEOL EX-230. Fig. 1(a) shows the OM micrograph of the Gd-modified as-forged sample. The mean grain size is about 50 μm, which is viewed from the transverse section. After annealing in the single α-phase region, the recrystallized microstructure with equiaxed grains can be seen. During cooling, the coarse α phase transforms into a fully lamellar structure which is composed of alternating laths of α2 and γ phases. The average grain sizes of the alloy gradually increase from 50 μm to 250 μm with annealing time up to 2 h (Fig. 1(b)). For comparison, we also investigated the grain evolution of Gd-free alloy. The grain size of the as-forged sample is about 80 μm, similar to that of Gd modified sample. However, after 2 h annealing, the grain size increases to over 1000 μm (Fig. 1(c)), which is four times larger than that of the 2 h annealed Gd-modified alloy. The grain sizes of the Gd-modified and Gd-free samples as a function of annealing time at 1340 °C are plotted in Fig. 1(d). It is obviously that the Gd addition can effectively prevent grain growth. In addition, the grain refinement effect of Gd addition for the as-cast alloys was also investigated. The grain sizes of the as-cast Gd-modified and Gd-free ingots from their columnar and equiaxed zones were measured. The Gd-modified sample has an average grain size of about 1000 μm in columnar zones and 200 μm in equiaxed zones while they are about 2000 μm and 300 μm, respectively, in the Gd-free ingot [18].

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To verify the effect of Gd on the retarding effect on grain growth during high temperature annealing in the single α-phase region, the Gd distribution in the TiAl alloy was firstly investigated by SEM. From the BSE images of the samples annealed at 1340 °C for different periods (Fig. 2), the Gd-rich precipitates with bright contrast can be clearly observed, which randomly distribute in the matrix in the as-forged sample (Fig. 2(a)). After annealing for 0.5 h, the size and distribution of these particles were almost no change. However, with the continuous increase of annealing time up to 2 h, the number of the Gd-rich precipitates gradually decreases with slight particle coarsening. Importantly, the Gd-rich precipitates tend to segregate along GBs to form precipitate-decorated GBs with the extension of annealing time (Fig. 2 (b)), which provides direct evidence that the slow grain growth of the Gd-modified TiAl alloy is associated with the GB precipitation of Gd-rich phases by dragging the GB migration. The chemical compositions of the Gd-rich precipitates were measured by SEM-EDS and STEM-EDS. From the survey of a large number of precipitates, we noticed that there are two kinds of Gd-rich phases. Particularly, STEM-EDS with a high spatial resolution can effectively minimize the influence from surrounding TiAl matrix. Most precipitates in the TiAl alloy are a Gd-Al intermetallic phase with a composition of 66.3 ± 3.1 at.% Al and 33.7 ± 2.2 at.% Gd and the other phase is a Gd oxide containing 59.8 ± 4.8 at.% O and 40.2 ± 2.1 at.% Gd. STEM-EDS suggests that the Gd oxide does not contain Ti, which is different with the possible phase Gd2TiO5 in the as-cast Gd-modified TiAl alloy suggested by Li and co-workers before [9]. Combined with the atomic scale structure characterization detailed below, the two Gd-rich phases are identified as Al2Gd and Gd2O3, respectively. Although the Gd-rich phases revealed by SEM often have a size in the order of micrometers, HAADF-STEM micrographs show that there are a large number of Gd-rich precipitates with a very small size of about several to tens of nanometers, about three orders of magnitude

Fig. 1. OM micrographs of the TiAl alloys after heat treatment at 1340 °C for different time: (a), (b) 0 h, and 2 h for Gd-modified samples, (c) 2 h for Gd-free sample and (d) the relationship between average grain size and annealing time for the two samples.

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Fig. 2. BSE images of the Gd-modified samples after annealing at 1340 °C for different time: (a) 0 h, (b) 2 h.

smaller than those imaged by SEM, in the annealed samples (Fig. 3(a)), which are invisible in the SEM images. The ultra-fine precipitates are believed to play the key role in microstructure refinement of Gd-modified TiAl. Similar to these coarse precipitates imaged by SEM, the nanosized precipitates tend to appear at GBs and lamellar interfaces (Fig. 3(b)),

indicating that the precipitates are not only functional in the grain refinement but also lamellar structure. The interaction between the precipitates and the formation of lamellar structure is evidenced by the enlarged HAADF-STEM image (in set of Fig. 3(b)), in which the lamellar interface cannot keep in a straight line after passing the nanoparticle.

Fig. 3. HAADF-STEM images of the precipitates in the Gd-modified sample annealed at 1340 °C for 2 h. (a) Low-magnification image; (b) Zoom-in image and the inset displays a nanoprecipitate at lamellar interface indicating the lamellar interface cannot keep in a straight line after passing it; (c) and (d) EDS mapping analysis of the nanoparticles marked with a rectangle and a circle in (b), respectively.

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Apparently, the interface shift around the precipitate leads to the variation of the lath interspacings. The interaction between the lamellar interfaces and nanoprecipitates also indicates that the precipitation of the nanoprecipitates takes place prior to the eutectoid reaction, mostly likely during cooling in the α phase regime. The nano-sized precipitates at GBs and lamellar interfaces were characterized by STEM-EDS (Fig. 3(c–d)) and have the same compositions as the coarse Al2Gd and Gd2O3 phases. Unlike the coarse Al2Gd and Gd2O3 precipitates which have obvious difference in morphology, the nano-sized precipitates of two phases have a very similar shape of being slightly elongated spherical particles. As confirmed by STEM-EDS, the Gd2O3 particles have a size ranging from 20 to 60 nm while the size of the Al2Gd precipitates is usually smaller than 20 nm. Since Gd has the largest atomic mass in the alloy, it can be easily distinguished from the HAADF-STEM images because of relatively bright contrast. Fig. 4(a) shows an Al2Gd particle embedding in a γ-TiAl lath. From the atomic image of the precipitate (Fig. 4(b)), the interplanar spacings of (222) and (400) are measured to be 0.223 nm and 0.199 nm, respectively. And thus the lattice constant is determined as 7.960 Å, which is well consistent with the standard database of C15 Al2Gd with the space group of Fd-3 m (227) and a lattice constant of 7.899 Å (JCPDS-PDF Card No. 28-0021(2004)). Therefore, the Al-Gd precipitates are determined to have a cubic MgCu2-type structure, i.e. the

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C15 Laves phase, which is, in fact, consistent with the stoichiometry of Al2Gd. Along the [110]γ orientation, the precipitate shows a regular shape with well-defined facets on {004} and {222} planes. It can be seen that the particle is slightly elongated with flat edges on the slow growth direction of [004] and the fast growth directions of [-222] and [2-22]. Particularly, on the fast growth {222} planes, atomic steps and even a single-atom-column terrace can be seen (Fig. 4(b)), suggesting that the {222} of the nanoprecipitate are kinetically unstable for fast growth. Moreover, the precipitate has a certain orientation relation with the γ-TiAl matrix, which is shown in the zoom-in HAADF-STEM image (Fig. 4(b)). According to the fast Fourier transform (FFT) pattern in which the orientations of both the γ-TiAl matrix and Al2Gd precipitates are indexed (the inset of Fig. 4(b)), the orientation relation can be described as: 〈022〉Al2Gd//〈110〉 γ, {2-22}Al2Gd//{002}γ with the deviation angle of about 4°. The crystallographic relationship seems to inherit from the precipitation of Al2Gd with the parent α-TiAl phase before the eutectoid reaction. The deviation may be caused by lattice strains and lattice distortions during the phase transformation from αTiAl to γ-TiAl and α2-Ti3Al. It is worth noting that atomic segregation of Gd atoms at GBs and lamellar interfaces cannot be detected by HAADF-STEM although extensive GB and interface precipitation of Gdrich phases takes place during high-temperature annealing and subsequent cooling.

Fig. 4. (a) HAADF-STEM image along 〈011〉γ zone axis of an Al2Gd particle inside the grain; (b) zoom-in image of the left bottom part in (a) and inserted FFT pattern; (c) HAADF-STEM image of a Gd2O3 particle at the lamella interface, the surface step is composed of many small facets along (004) and (222) planes, and the inset showing a low-magnification bright-field TEM image of this precipitate at the lamella interface; (d) enlarged image of (c) and inserted FFT pattern.

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The atomic structure of the Gd2O3 phase was also characterized by HAADF-STEM (Fig. 4(c,d)). From the zoom-in atomic images in Fig. 4(d), the interplanar spacings of (004) and (222) of the oxide are measured to be 0.269 nm and 0.309 nm, respectively, and thus the lattice constant is determined as 10.760 Å, which agree with those of Gd2O3 with the space group of Ia-3 (206) and a lattice constant of 10.813 Å (JCPDS-PDF Card No. 12-0797(2004)). Most oxide nanoparticles locate at GBs and interfaces and have a size obviously larger than that of Al2Gd (Fig. 3(a,b)). The inset in Fig. 4(c) is a bright-field TEM image showing an oxide precipitate at a GB. From the contrast variation around the nanoparticle, it can be identified that the particle resides at a GB where a γ-TiAl/α2-Ti3Al interface also ends at the precipitate. We have systematically tilted the Gd2O3 nanoparticle to different orientations. When a Gd2O3 nanoparticle is aligned to a low index zone orientation and can be imaged with clear atomic columns, the TiAl matrix is always out of a zone orientation (Fig. 4(c)). Thus, unlike the Al2Gd precipitates, the Gd2O3 nanoparticles do not have a well-defined crystallographic relationship with the γ-TiAl and α2-Ti3Al matrix. From the atomic image and corresponding FFT pattern (Fig. 4(d)), the oxide is determined to have an ordered face-centered cubic structure, which is in agreement with the known crystal structure of Gd2O3. From the atomic-scale image (Fig. 4(c)), it can be found that the spherical shape of the particle is formed by a short steps composed of small facets along {004} and {222} planes. A large facet and a buffer layer around the oxide cannot be seen. The incoherent interface between the Gd2O3 precipitates and the TiAl matrix could be the key reason leading to a relatively higher growth rate than that of Al2Gd. The solubility of Gd in TiAl is very low although the value changes with the compositions of parent alloys and annealing temperatures [19]. In this study, the significant microstructure refinement has been observed in the TiAl alloy containing only 0.15 at.% Gd, indicating a high alloying efficiency of Gd in the microstructure modification of fully lamellar TiAl. According to the heat of mixing, Gd has a stronger bonding with Al than Ti and thus it is reasonable that a Gd-Al phase precipitates preferentially from the TiAl alloy. Based on the binary Al-Gd phase diagram, several high-temperature intermetallic phases with different melting points can be formed in the Al-Gd system, including Gd2Al (950 °C), Gd3Al2 (980 °C), GdAl (1075 °C), GdAl3 (1125 °C) and Al2Gd (1525 °C) [20]. However, only the Al2Gd phase can survive at the annealing temperature of 1340 °C, which is consistent with the TEM characterization. This consistency also confirms that the precipitation of Al2Gd, in particularly the coarse precipitates along GBs, takes place during high temperature annealing in the single α-phase region, which thus leads to the effective grain refinement. For the nano-sized precipitates in the interior grains and lamellar interfaces, the precipitation may occur during air cooling before the eutectoid reaction because of the imperfect crystallographic relation between the nano-precipitate and TiAl matrix as well as the strong interaction between the precipitates and lamellar interfaces. Besides the Al2Gd intermetallic compound, the Gd2O3 oxide, particularly nano-sized precipitates, can also be detected in the alloy. REEs have been used extensively to purify alloys because of their high affinity with oxygen and, in principle, would preferentially form oxides in an alloy containing oxygen. Since Ti and Al also have strong affiliation with oxygen, impurity oxygen in TiAl alloy is unavoidable and is known to affect the mechanical properties of TiAl alloys [21]. In the Gd-modified TiAl alloy, Gd apparently has the function in the purification of the TiAl by forming Gd2O3 because of the high melting temperature (2330 °C) of Gd2O3 and lower solubility of Gd and O in the TiAl melt. As a high-temperature solid phase, the oxide particles usually have a large size and distribute in the finally solidified regions, such as

GBs, and correspond to the coarse oxide precipitates (Fig. 2). Different from the coarse Gd2O3 phase from TiAl melt, the HAADF-STEM reveals the formation of nano-size Gd2O3 precipitates from α-TiAl solid solution. Importantly, the oxide nanoprecipitates at GBs and lamellar interfaces play the same role as the designed Al2Gd phase in the microstructure refinement. This finding has an important implication in utilizing oxide nanoprecipitates for microstructure refinements and mechanical property enhancement of TiAl alloys and may open a door for the development of oxide dispersion strengthened (ODS) TiAl alloys by casting through alloy design and optimization of thermal processing. We have not detected segregated Gd at GBs or lamellar interfaces by using HAADF-STEM. The grain refinement effects seem to be solely from Gd2O3 and Al2Gd precipitates. Since the melting temperature of Gd2O3 is significantly higher than the annealing temperature, the coarse oxides appears to be stable during the solid solution treatment at 1340 °C and the grain refinement is mainly by pinning GBs. For the Al2Gd intermetallic phase, visible morphology changes of the coarse precipitates during annealing suggest that the intermetallic phase is unstable during the solution treatment and experiences Ostwald ripening. The obvious GB precipitation, which may be driven by the high interface energy, evidently prevents the grain coarsening of the TiAl alloy by dragging the GB motion. In summary, the effect of Gd addition on the microstructure refinement of a fully lamellar TiAl alloy was systematically studied. Two types of Gd-contained precipitates, Gd2O3 and Al2Gd, with sizes ranging from several nanometers to micrometers were detected at GBs and lamellar interfaces, suggesting that the microstructure refinements are from both oxide and intermetallic precipitates. In particularly, the refinement effect from the oxides has an important implication in developing ODS TiAl alloys. Acknowledgments The authors are grateful for financial support from the National Natural Science Foundation of China (Grant no. 51271112, 11327902, 51271113), MOST 973 of China (Grant No. 2015CB856800). References [1] C.T. Liu, P.J. Maziasz, Intermetallics 6 (1998) 653. [2] K. Kothari, R. Radhakrishnan, N.M. Wereley, Prog. Aerosp. Sci. 55 (2012) 1. [3] W. Wallgram, T. Schmolzer, L.M. Cha, G. Das, V. Guther, H. Clemens, J. Mater. Res. 100 (2009) 1021. [4] J.N. Wang, J. Zhu, J.S. Wu, X.W. Du, Acta Mater. 50 (2002) 1307. [5] T. Klein, B. Rashkova, D. Holec, H. Clemens, S. Mayer, Acta Mater. 110 (2016) 236. [6] K. Xia, W. Li, C. Liu, Scr. Mater. 41 (1999) 67. [7] Y. Wu, S.K. Hwang, Acta Mater. 50 (2002) 1479. [8] F.S. Sun, C.X. Cao, S.E. Kim, Y.T. Lee, M.G. Yan, Scr. Mater. 44 (2001) 2775. [9] W. Li, B. Inkson, Z. Horita, K. Xia, Intermetallics 8 (2000) 519. [10] K. Xia, X. Wu, D. Song, Acta Mater. 52 (2004) 841. [11] W. Li, K. Xia, Mater. Sci. Eng. A 329–331 (2002) 430. [12] L. Chen, X.C. Ma, M. Jin, J. Wang, H. Long, Metall. Mater. Trans. A A47 (2016) 33. [13] X. Hu, F. Jiang, F. Ai, H. Yan, J. Alloys Compd. 538 (2012) 21. [14] H. Pan, Y. Ren, H. Fu, H. Zhao, L. Wang, X. Meng, G. Qin, J. Alloys Compd. 663 (2016) 321. [15] S. Mitao, L.A. Bendersky, Acta Mater. 45 (1997) 4475. [16] E. Schwaighofer, P. Staron, B. Rashkova, A. Stark, N. Schell, H. Clemens, S. Mayer, Acta Mater. 77 (2014) 360. [17] A. Denquin, S. Naka, Acta Mater. 44 (1996) 343. [18] X. Chang, J.Y. Si, P.B. Han, J. Zhang, J. Iron Steel Res. 21 (2009) 47 (in Chinese edition). [19] C. Liu, W. Li, K. Xia, in: T. Chandra, T. Sakai (Eds.), Proc. Inter. Conf. on Materials, Processing at High Temperature, Wollongong, NSW, Australia, Vol. 2, 1997, pp. 1575–1581. [20] H. Okamoto, Phase Diagrams of Dilute Binary Alloys, Materials Park, ASM International, 2002. [21] M. Morris, Intermetallics 4 (1996) 417.