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Microstructural evolution and high-temperature oxidation mechanisms of a titanium aluminide based alloy
Accepted Manuscript Microstructural evolution and high-temperature oxidation mechanisms of a titanium aluminide based alloy S.J. Qu, S.Q. Tang, A.H. Feng, C. Feng, J. Shen, D.L. Chen PII:
S1359-6454(18)30114-9
DOI:
10.1016/j.actamat.2018.02.013
Reference:
AM 14369
To appear in:
Acta Materialia
Received Date: 26 November 2017 Revised Date:
28 January 2018
Accepted Date: 5 February 2018
Please cite this article as: S.J. Qu, S.Q. Tang, A.H. Feng, C. Feng, J. Shen, D.L. Chen, Microstructural evolution and high-temperature oxidation mechanisms of a titanium aluminide based alloy, Acta Materialia (2018), doi: 10.1016/j.actamat.2018.02.013. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
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ACCEPTED MANUSCRIPT
ACCEPTED MANUSCRIPT Microstructural evolution and high-temperature oxidation mechanisms of a titanium aluminide based alloy
S.J. Qu1*, S.Q. Tang1, A.H. Feng1, C. Feng1, J. Shen1, D.L. Chen2∗ School of Materials Science and Engineering, Tongji University, Shanghai 201804, China 2
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1
Department of Mechanical and Industrial Engineering, Ryerson University, Toronto,
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Ontario M5B 2K3, Canada
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Abstract: Oxidation resistance of titanium aluminide (TiAl) based alloys is a fundamental aspect for the high-temperature structural applications such as in the advanced hypersonic aircraft engines and gas turbines. The aim of this study was to identify oxidation kinetics and mechanisms through detailed microstructural characterization of a newly-developed
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Ti-44Al-4Nb-1.5Cr-0.5Mo-0.1B-0.1Y alloy via focused ion beam (FIB), transmission electron microscopy (TEM), X-ray diffraction (XRD), electron probe microanalysis (EPMA),
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scanning transmission electron microscopy (STEM), along with density functional theory (DFT) calculations. The alloy consisting mainly of γ-TiAl/α2-Ti3Al lamellar structure
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exhibited a superior oxidation resistance at 700°C, and followed parabolic oxidation kinetics at 800°C and 900°C. The observed multi-layered scale structure consisted of TiO2, Al2O3-rich, Al2O3+TiO2,
H-Ti2AlN+Al2O3+α2-Ti3Al,
Z-Ti5Al3O2+AlNb2+Laves-(Ti,Nb)Cr2,
and
H-Ti2AlN/α2-Ti3Al lamellae from the outside to inside after high-temperature oxidation. The γ-TiAl/α2-Ti3Al lamellae near the scale/substrate interface were first transformed into
∗
Corresponding author. Tel.: +86-21-6958-1009; E-mail address: [email protected] (S.J. Qu); and Tel: (416) 979-5000 ext. 6487; Fax: (416) 979-5265; Email: [email protected] (D.L. Chen). 1
ACCEPTED MANUSCRIPT H-Ti2AlN/α2-Ti3Al lamellae, with orientation relationships identified as (0001)α2 //(0001)Ti2AlN , (10 1 0)α 2 //(10 1 0)Ti 2AlN and [12 10]α 2 //[1210]Ti 2AlN . The H-Ti2AlN/α2-Ti3Al lamellae were
then transformed into a metastable Z-Ti5Al3O2 phase at the scale/substrate interface. The
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Z-phase was decomposed to Ti3Al and Al2O3 as the scale/substrate interface moved inwardly. Ti3Al reacted further with oxygen and nitrogen to form Ti2AlN, which was finally oxidized to form TiO2 and α-Al2O3. A Nb-rich layer was present beneath the scale along with the
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formation of AlNb2 and Laves phase, and the doping effect of Nb to suppress the diffusion of
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oxygen occurred mainly in the TiO2+Al2O3 compound layer. The results obtained in this study would pave the way for the development of advanced oxidation-resistant TiAl-based materials for high-temperature applications.
1. Introduction
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Keywords: TiAl-based alloys; oxidation kinetics; oxidation mechanisms; EPMA; TEM.
As a relatively new class of promising materials for high-temperature structural
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applications such as in the advanced hypersonic aircraft engines and gas turbines,
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intermetallic titanium aluminide (TiAl) based alloys have recently drawn considerable attention due to their attractive characteristics, including lightweighting (with a low density of ρ≈3.9-4.2 g/cm3), high specific strength, excellent modulus retention, and superior creep resistance [1-6]. This is indeed motivated by the successful application of a TiAl-based alloy in General Electric’s high-thrust GEnex jet engine for powering Boeing 747-8 and 787 Dreamliner [2,3,7,8]. TiAl-based alloys in polycrystalline forms have been increasingly utilized to replace nickel-based superalloys (with a density of ρ≈8.5 g/cm3) in the temperature 2
ACCEPTED MANUSCRIPT range of 650-750°C, with the benefit of ~50% weight reduction [2,3,9,10]. However, the widespread commercial applications of these alloys are still limited due to the low room-temperature ductility and some challenges in manufacturing [1,11], and especially
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unsatisfactory oxidation resistance above ~750°C [12,13]. Because of a similar activity of Ti and Al in the severe environment of turbine engines, a protective Al2O3 layer could not be properly formed on the surface of TiAl-based alloys after
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high-temperature oxidation [14]. Generally, nonprotective multi-layered oxide scales
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composed of TiO2/Al2O3/TiO2+Al2O3 occurred after oxidation of TiAl, since γ-TiAl cannot be directly oxidized to TiO2 and Al2O3 [15-17]. During oxidation titanium nitrides (TiN and Ti2AlN) along with a metastable cubic Z-phase (Ti5Al3O2) (or X-phase) have been observed at the interface of oxide scale and substrate via TEM [18-20]. Copland et al. [21] reported
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that in binary alloys containing γ-TiAl and α2-Ti3Al phases, metastable Z-phase first occurred from the oxidation of γ-TiAl and eventually decomposed to α2-Ti3Al (rich in oxygen) and Al2O3 when oxidation experiments were conducted at 1000°C in a high-purity oxygen
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environment. Dettenwanger et al. [22] observed an Al-depletion layer after oxidation of a
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binary Ti-50Al alloy at 900°C for 1 h in air, where the Z-phase was also the first phase formed. In the subsequent oxidation, α2-Ti3Al will form at the interface of Z-phase and γ-TiAl. The Z-phase with a higher Vickers hardness may induce cracks in the subsurface zone and Al2O3 layer [21,22]. In addition, the Z-phase could destroy the protective Al2O3 layer after its decomposition to α2-Ti3Al and Al2O3 [21]. Titanium nitrides were observed only at the substrate/scale interface where the ratio of p O2 p N 2 is low, because the formation of the oxides leads to a reduction of p O 2 [15]. Nitrides would be oxidized to TiO2 after long 3
ACCEPTED MANUSCRIPT exposure and the released nitrogen can result in the formation of a new nitride layer. This repeated process thus impedes the formation of a continuous Al2O3 layer [22]. The orientation relationship between the nitrides in the nitride layer was observed to be (111)TiN || (0001)Ti AlN 2
and [1 11]TiN || [1120]Ti AlN . High-density stacking faults in the TiN phase and Al diffusion from
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2
the substrate to the outer layer lead to the formation of the Ti2AlN phase at the bottom of the TiN phase [23]. However, it is unclear how the Z-phase evolves into nitrides and if the
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lamellar structure below the Z-phase would change in the oxidation process.
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To hinder the diffusion of O and N from air to substrate, several measures were taken [24], such as coating technologies [25-27] and surface treatment [28-31]. Adding moderate ternary elements, such as Nb, W, Si, Mo, Cr and Y in TiAl-based alloys, improves not only the mechanical properties but also the oxidation resistance at high temperatures [32-37].
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Dopant elements with a valence higher than Ti4+, such as Nb5+, would cause a reduction of O vacancy concentration in the TiO2 lattice in order to maintain electrical neutrality. Thus the growth of TiO2 is blocked at a lower oxygen diffusion rate [38]. A continuous sub-scale
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Ti(Cr,Al)2 Laves phase could form if a certain amount (7~10 at.%) of Cr was added to
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TiAl-based alloys [39,40]. The addition of element Y results in a reduction of weight gain and an improvement of adhesion between oxide scale and substrate [34,41]. Although the role of the added elements was discussed, their influence on the microstructural evolution of Z-phase,
α2-Ti3Al and substrate was rarely identified. Besides, it is unknown which phase in the γ-TiAl/α2-Ti3Al lamellar structure would change during high-temperature oxidation. The purpose of this study was, therefore, to identify the microstructural evolution from
γ-TiAl/α2-Ti3Al to Al2O3 (corundum) and TiO2 (rutile) and fundamental oxidation 4
ACCEPTED MANUSCRIPT mechanisms via detailed characterization of the whole oxide scale after oxidation of a newly-developed TiAlNbCr alloy at 900°C for 100 h via SEM, STEM, EPMA, FIB and TEM,
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along with relevant thermodynamic calculations involving density functional theory (DFT).
2.1. Specimen preparation material
with
a
nominal
chemical
composition
of
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The
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2. Experimental procedures
Ti-44Al-4Nb-1.5Cr-0.5Mo-0.1B-0.1Y (at.%) (hereafter referred to as TiAlNbCr alloy) was prepared by double vacuum consumable arc melting technique. For oxidation tests, the specimens with approximate dimensions of 10×10×1 mm3 were cut from the alloy ingots by
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wire EDM. The surfaces were ground using SiC sandpapers up to grit No. 1200, and then cleaned ultrasonically in acetone for 15 min. Finally, the initial size and weight of each
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specimen were measured for calculating weight gain later.
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2.2. Oxidation tests
Isothermal oxidation tests were carried out at 700, 800 and 900°C in a muffle furnace in static laboratory air for 100 h. A hole on the door with a diameter of 10 mm was drilled to ensure sufficient oxygen in the furnace chamber. After a specified temperature has reached, multiple specimens were put in the furnace and then individual specimens were taken out after 1, 3, 6, 12, 24, 36, 50, 62, 74, 86 and 100 h, respectively. Their weights were measured by an electronic balance with a resolution of 0.1 mg to calculate the weight gain ∆M. 5
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2.3. Microstructural characterization Original microstructures of the test material were examined by field emission
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environmental scanning electron microscope (FE-ESEM, FEI Quanta 250 FEG) equipped with energy-dispersive X-ray spectroscopy (EDS, Apollo SDD10), and transmission electron microscopy (TEM, FEI Talos 200X). X-ray diffraction (XRD, Rigaku D/Max-2550) with Cu
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Kα radiation (λ=1.5418Å) was used to identify the phases in the oxide scale at 50 kV and 200
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mA with a diffraction angle (2θ) from 10° to 100° with a step size of 0.02° and 1 s in each step. Elemental distribution in different regions on the cross-section of the oxide scale was characterized by using the electron probe microanalyzer (EPMA, Shimadzu 1720) with a resolution of 1 µm and secondary-electron image resolution of 6 nm using a beam current of
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10 nA, and scanning transmission electron microscopy (STEM, FEI Talos 200X) with EDS. In the EPMA the standard samples used for calibrating the elemental distributions of titanium, aluminum, oxygen, nitrogen, niobium, chromium, molybdenum, boron and yttrium were pure
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titanium, Al2O3, BN, pure niobium, pure chromium, pure molybdenum, pure boron and
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YP4O12, respectively. Phase constitution at the interface of scale and substrate was determined by focused ion beam (FIB, FEI NanoLab 600i) and TEM.
2.4 Computational details
Thermodynamic calculations were performed using Materials Studio DMol3 version 6.0.0 (Accelrys Software Inc.), a high-quality quantum mechanics program [42]. DMol3 uses density functional theory (DFT) with a numerical radial function [43] basis set to calculate 6
ACCEPTED MANUSCRIPT the electronic properties of molecules, surfaces and crystalline solid materials [44] from the first-principles. Geometrical optimizations and frequency calculations were carried out with the GGA-BLYP function in conjunction with a double numerical plus d-functions (DND)
3. Results
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3.1. Original microstructures of TiAlNbCr alloy
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basis set to calculate free energy (G) as a function of temperature.
Back-scattered electron (BSE) SEM micrograph and TEM bright field (BF) microstructure of as-cast TiAlNbCr alloy are shown in Fig. 1(a) and (b). The microstructure consisted of grains (or colonies) containing characteristic (γ-TiAl + α2-Ti3Al) lamellae, where
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the bright B2-phase (ordered β-Ti) and dark γ-phase were present at the grain boundaries (Fig. 1(a)). The representative γ+α2 lamellar structure could be better seen from the TEM image
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along with the selected area diffraction (SAD) patterns in Fig. 1(b).
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3.2. Isothermal oxidation kinetics
Curves of isothermal-oxidation kinetics of TiAlNbCr alloy at 700, 800 and 900°C are
shown in Fig. 2. The obtained mass gain of this alloy after 100 h at 700, 800 and 900°C was about 0.0033, 0.4482 and 2.3777 mg/cm2, respectively. Thus this alloy exhibited a superior oxidation resistance at 700°C, with respect to the situation at 800 and 900°C. To identify which law of oxidation kinetics is followed, the obtained experimental data was fitted using the following equation, 7
ACCEPTED MANUSCRIPT ∆M n = k n t ,
(1)
where ∆M represents weight gain per unit area (mg/cm2), n is an oxidation exponent (n=1, linear relationship; n=2, parabolic relationship), kn is a rate constant (mgn/cm2n h) and t is
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oxidation time (h). The fitting curves via Origin software were plotted in Fig. 2, and the oxidation exponent (n) at 800 and 900°C was obtained to be 1.83 and 1.80, respectively. Hence the oxidation at 800 and 900°C obeyed a parabolic growth kinetics (n≈2) from 0 to
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100 h. Under this circumstance the value of kn would reflect the oxidation rate. Compared
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with kn=2.44×10-3 mg2/cm4 h at 800°C, the oxidation rate constant at 900°C was much higher (kn = 4.99×10-2 mg2/cm4 h). It suggests that this alloy exhibited a poor oxidation resistance at 900°C. Therefore, the subsequent study will be mainly focused on the oxidation of this alloy
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at 900°C.
3.3. XRD results
X-ray diffraction patterns of oxide scales in TiAlNbCr alloy are shown in Fig. 3. It is
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seen that after oxidation at 700°C for 100 h, except the γ, α2 and B2 in the substrate, oxides
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could barely be detected, while the peaks of oxides (rutile TiO2 and corundum α-Al2O3) and nitrides (TiN and Ti2AlN) were obvious at 800°C. When the temperature increased to 900°C, the peaks of rutile and corundum were further elevated, implying an increasing extent of oxidation. Furthermore, while the peaks of Ti2AlN became stronger, the peaks of TiN disappeared at 900°C.
3.4. Cross-sectional observations 8
ACCEPTED MANUSCRIPT Fig. 4(a) presents the SEM-BSE micrograph of oxide scale after 100 h of oxidation at 900°C. According to the elemental distribution maps of the cross-section of the oxide scale detected by EPMA shown in Fig. 4(b), it is seen that the outermost grey layer I consisted of
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columnar particles, and the adjacent black layer II was enriched in O and Al. Al and Ti coexisted in a mixed grey and black layer III. Combined with the XRD patterns shown in Fig. 3, the scale structure formed was identified to be in the order of TiO2 layer (I)/Al2O3-rich
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layer (II)/mixed Al2O3 +TiO2 layer (III) from the outside to inside.
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To reveal the chemical composition quantitatively in various areas of the cross-sectional scales, EPMA point microanalyses were conducted, as indicated in Fig. 4(a). The obtained local chemical compositions of points 1-16 are summarized in Table 1. Nb was rich in the white layer (about 3~4 µm) beneath the interface of oxide scale and substrate as indicated by
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points 10 and 11 in Fig. 4(a) and Nb distribution map in Fig. 4(b). In the oxide scale, Nb was mainly distributed in the layer III of TiO2 + Al2O3 mixture (points 5-9), but absent in the outer TiO2 layer I (points 1 and 2) and mid Al2O3-rich layer II (points 3 and 4). According to the
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chemical compositions at points 8 and 9, a small amount of nitrides was presumably formed
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at the scale/substrate interface because a sharp increase of N content was also detected (Table 1), which will be confirmed via subsequent TEM examinations. Besides, N and O with a concentration of 5~9% permeated into the lamellar colonies in the vicinity of scale/substrate interface, as indicated by points 12 and 13 (Table 1), which might cause a potential reduction of ductility of the alloy. Cr was enriched along with Mo and Nb in the white areas marked by points 14, 15 and 16 at the grain boundaries close to the scale, which corresponded to the presence of B2 phase in the substrate. B2 phase could 9
ACCEPTED MANUSCRIPT promote the formation of Al2O3, since the diffusion of Al was faster in B2 phase than in
γ-TiAl [45].
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3.5 TEM and STEM investigations The TEM/STEM sample of TiAlNbCr alloy after oxidation at 900°C for 100 h was carefully positioned at the scale/substrate interface, which was prepared via a
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focused-ion-beam (FIB) technique, as shown in Fig. 5. Fig. 6(a) and (b) show a TEM-BF
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image and a high-angle annular dark-field (HAADF) image at the scale/substrate interface, respectively. The lamellar structure at the bottom of Fig. 6(a) and (b) reflected the alloy substrate, and the equiaxed ultra-fine crystallites located above the lamellae were the microstructure of oxide scale. The elemental distribution in the entire area was very
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non-uniform as seen from Fig. 6(c)-(k), where the STEM-EDS elemental mapping of Ti, Al, O, N, Nb, Cr, Mo, B and Y was shown. The oxide scale mainly contained Ti, Al and O, and beneath the oxides there existed a distinct N-rich layer (Fig. 6(f)), which superimposed nicely
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with the location of Ti (Fig. 6(c)) and was in good agreement with the EPMA microanalysis
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results of points 8 and 9 in Fig. 4(a) and Table 1. Similarly, the positions of Al and O were also well overlapped, as seen from Fig. 6(d) and (e). The white particles present at the scale/substrate interface as seen from Fig. 6(b) were rich in Nb, Cr and Mo (Fig. 6(g)-(i)). More detailed TEM and high-resolution transmission electron microscopy (HRTEM) examinations of the interfacial area between the oxide scale and lamellar substrate were performed, as shown in Fig. 7. The oval area marked by “1” at the bottom-left of Fig. 7(a) had an obviously different morphology compared with the adjoining inner lamella and outer 10
ACCEPTED MANUSCRIPT ultra-fine oxide particles. The SAD pattern shown in Fig. 7(b) indicated that this oval block consisted of a simple cubic Z-phase (Ti5Al3O2), which exhibited a significantly higher Vickers hardness than base metal [22]. Another Z-phase was spotted in the area marked by
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“2”, which was far from area “1”. Areas “3” and “4” were identified as hexagonal H-phase (Ti2AlN) [46-48] through the SAD patterns (Fig. 7(d) and (e)). These discontinuous nitrides were present only at the scale/substrate interface, as also shown in Fig. 6(f). A characteristic
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microstructure of islands in area “5” consisted of alumina (α-Al2O3) formed besides the
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nitrides at the scale/substrate interface, corresponding well to Al and O mapping in Fig. 6(d) and (e). This suggests that the growth of α-Al2O3 oxides was interrupted by nitrides, making it difficult or impossible to form a dense protective layer, as also noted by Lu et al. [49]. The porous oxide layer outside the interface was mainly composed of a great deal of ultra-fine
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equiaxed rutile TiO2 such as areas “7” and “8” and a small amount of alumina (α-Al2O3, as indicated by area “6”). The area close to Z-phase marked by “9” was identified to be α2-Ti3Al and the other equiaxed α2 was only observed in area “10” in Fig. 7(k), which is a magnified
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image of zone “A” boxed in Fig. 7(a).
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Furthermore, the white particles at the scale/substrate interface shown in Fig. 8(a) were rich in Nb, Cr and Mo according to the STEM-EDS results in Fig. 6(g)-(i). Fig. 8(b) and (c) show higher magnification HAADF images of zones “A” and “B” in Fig. 8(a), where point “ Ι ” in Fig. 8(b) was enriched in Nb, i.e., the phase of AlNb2 as revealed by the SAD results (Fig. 8(d)), and point “ ΙΙ ” rich in Nb and Cr in Fig. 8(c) was identified as Laves phase (Nb, Ti)Cr2 (Fig. 8(e)). While the lamellar structure in the neighborhood of scale/substrate interface in Fig. 9(a) 11
ACCEPTED MANUSCRIPT looked seemingly like that in Fig. 1(b), the SAD results revealed that the outermost layer marked “A” of this specific lamella was neither α2-Ti3Al nor γ-TiAl, but the H-phase (Ti2AlN). SAD pattern from point “B” (i.e., the interface of two lamellae) showed that the
between
these
two
phases
were
(0001)α2 //(0001)Ti2AlN ,
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lamella adjacent to “A” still remained to be α2-Ti3Al, and the orientation relationships (10 1 0) α2 //(10 1 0) Ti2AlN
and
[1210]α 2 //[1210]Ti 2AlN . The inner layers “C” and “D” corresponded to H-Ti2AlN and α2-Ti3Al,
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respectively, and point “E” at some distance from the scale/substrate interface still had the
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same SAD pattern as point “B”. As a result, there was no γ-TiAl detected in the visible lamellar colony near the scale/substrate interface, while the α2-Ti3Al lamella remained unchanged after oxidation. This suggests that the original γ-TiAl/α2-Ti3Al lamellar structure adjacent to the scale was replaced by the H-Ti2AlN/α2-Ti3Al lamellar structure after oxidation
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4. Discussion
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at 900°C for 100 h.
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4.1 Multi-layered microstructural characteristics of oxide scale After isothermal oxidation in laboratory air at 900°C for 100 h, the oxide scale formed on the surface of TiAlNbCr alloy with a thickness of ~13 µm, which consisted of multiple layers of different oxides, nitride and intermetallic compounds. Based on the above examinations, the multi-layers from the outside to the inside were TiO2, Al2O3-rich, Al2O3+TiO2,
H-Ti2AlN+Al2O3+α2-Ti3Al,
Z-Ti5Al3O2+AlNb2+Laves-(Ti,Nb)Cr2,
H-Ti2AlN/α2-Ti3Al lamellae, and then alloy substrate (i.e., original γ-TiAl/α2-Ti3Al lamellar 12
ACCEPTED MANUSCRIPT structure). TiO2 and Al2O3 were the main oxidation products because they were thermodynamically stable in the Ti-Al-O-N system when oxygen was sufficiently offered [50]. TiO2 was a fast-growing oxide that did not provide long-term oxidation protection [51],
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consequently it occupied the largest scale volume. Discontinuous Z-phase, AlNb2 and Laves phase were present at the scale/substrate interface beneath the Ti2AlN-rich layer. The
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formation and decomposition of these phases will be discussed later.
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4.2 Oxidation kinetics
TiAlNbCr alloy exhibited a superior oxidation resistance at 700°C, as seen from Fig. 2, but weight gains became increasingly obvious at 800 and 900°C which basically obeyed a parabolic law. This suggests that the growth of oxide scale was controlled by diffusion
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processes because the parabolic law could be derived from Fick's first law of diffusion. During the parabolic oxidation of TiAl-based alloys, most of the scale formed by inward diffusion of oxygen [52]. TiO2 was the non-metal deficient n-type oxide with many oxygen
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vacancies and Ti4+ interstitial cations and thus became the main diffusion path [38]. The
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doping of Nb5+ was expected to decrease the oxygen vacancy concentration in order to maintain electroneutrality in TiO2 [53]. The relationship between diffusion coefficient D and the vacancy concentration Cv could be expressed by [54],
D = ACv F ,
(2)
where A is a crystal lattice-related parameter and F is the successful jump frequency of the atom. It is clear that oxygen diffusion would be suppressed when oxygen vacancies decreased. Jiang et al. [55] also reported that the oxidation resistance of TiAl-Nb increased with 13
ACCEPTED MANUSCRIPT increasing concentration of Nb from 0 to 20%. Therefore, the addition of Nb in the present alloy could decrease oxidation rate. That is, the observed white Nb-rich layer beneath the scale in Fig. 8(a) could be considered as a Nb-rich reservoir to block the inward diffusion of
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oxygen to a certain extent. However, further studies in this aspect are needed.
4.3 Formation of H-Ti2AlN/α2-Ti3Al lamellae
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It is of particular interest to observe that the H-Ti2AlN/α2-Ti3Al lamellar structure formed
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under the layer of Z-phase and Nb, Cr-rich phases (AlNb2 and Laves) to replace the original
γ-TiAl/α2-Ti3Al lamellae (Fig. 9). To understand this phenomenon, module DMol3 in the software of Material Studio 6.0 was used to assess thermodynamic possibility of phase transformation from γ-TiAl or α2-Ti3Al to H-Ti2AlN. Primitive crystal structures of γ-TiAl,
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α2-Ti3Al and H-Ti2AlN which are needed in the DFT calculations are shown in Fig. 10, with their lattice constants and atomic coordinates from ICSD database listed in Table 2. The dots in Fig. 11 represent the change of free energy as a function of temperature from 0-1000 K for
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γ-TiAl, α2-Ti3Al and H-Ti2AlN calculated via DMol3, and the corresponding curves were
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fitted according to the following equation, G = H 0 + aT ln T + bT 2 + c T + IT ,
(3)
where a, b, c, and I are all fitting parameters. The results obtained via the fitting are shown in Table 3 with an adjusted R-squared very close to 1. The Gibbs free energy change based on the following reactions,
4TiAl + 2 N → 2Ti 2AlN + 2Al ,
(4)
2Ti 3Al + 2 N → 2Ti 2 AlN + 2Ti ,
(5)
14
ACCEPTED MANUSCRIPT at 900°C are obtained to be -77.6 kJ/mol and 26.2 kJ/mol, respectively. This indicates that the transformation of γ-TiAl to H-Ti2AlN was thermodynamically possible, while the transformation of α2-Ti3Al to H-Ti2AlN was impossible. Therefore, there was a strong
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tendency of transformation from γ-TiAl/α2-Ti3Al lamellae to H-Ti2AlN/α2-Ti3Al lamellae due to the presence of thermodynamic driving force at 900°C. Indeed, such a transformation
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would be possible when the temperature was above ~420 K, as seen from Fig. 11.
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4.4 Formation and decomposition of Z-phase
The metastable Z-phase with a chemical formula of Ti5Al3O2 was observed at the scale frontier, i.e., at the interface of scale and H-Ti2AlN/α2-Ti3Al lamellae (Fig. 7). Dettenwanger et al. [22] and Copland et al. [21] reported that for binary Ti-50Al alloy and Ti-52Al
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single-phase γ-TiAl alloy, the Z-phase formed in the depletion layer stayed continuous after oxidation at 900°C for 100 h and at 1000°C for 50 h, respectively. According to Lu et al. [49] and Copland et al. [21], the Z-phase came from the oxygen absorption of substrate alloy, but
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our TEM observations showed that the Z-phase was present discontinuously, which was
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surrounded by H-Ti2AlN/α2-Ti3Al lamellae, AlNb2, α-Al2O3 and H-Ti2AlN as shown in Fig. 7(a) and Fig. 8(a). Then the only possible source of Z-phase would be related to the H-Ti2AlN/α2-Ti3Al lamellae rather than the substrate γ-TiAl/α2-Ti3Al lamellae, and the formation process could be expressed as follows (where the chemical formula of H-Ti2AlN/α2-Ti3Al
lamellae
could
be
re-written
as
[50Ti-44Al-4Nb-1.5Cr-0.5Mo-0.1B-0.1Y)]xN1-x in our material), [50Ti -44Al-4Nb-1.5Cr -0.5Mo-0.1B-0.1Y )]x N (1− x ) + 20 xO → 10 xTi5 Al3O2 + 14 xAl + 4 xNb + 1.5xCr+0.5xMo+0.1xB+0.1xY + (1-x)N . (6) 15
ACCEPTED MANUSCRIPT This suggests that there appeared an enrichment of Al, Nb and N after the formation of Z-phase. Then AlNb2 and Laves phase of (Ti, Nb)Cr2 would form close to the Z-phase as shown in Fig. 8(a) and these two phases would impede continuous Z-phase formation in the
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further oxidation [49]. Laves phase (Ti, Nb)Cr2 had a good resistance to oxidation [56], and thus exposure of Z-phase and H-Ti2AlN/α2-Ti3Al lamellae to the oxygen could be alleviated to some extent because of the presence of AlNb2 and Laves phase at the sub-interface. It is
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expected that the oxidation resistance of TiAl-based alloys could be greatly improved if the
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added Nb and Cr was sufficient to form a continuous layer consisting of AlNb2 and Laves phases. The N atom released according to Equ. (6) would continue to diffuse inwards to react with inner γ-TiAl to form H-Ti2AlN based on Equ. (4), further corroborating the formation of H-Ti2AlN/α2-Ti3Al lamellae, as discussed above. It should be noted that the Z-phase could
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exist in equilibrium with γ-TiAl, α2-Ti3Al and α-A12O3 at specific oxygen concentrations [57]; it could also decompose in the following oxidation process [21,49], (7)
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3Ti 5 Al3O 2 → 2Al2O3 + 5Ti 3Al .
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4.5 Formation and decomposition of H-Ti2AlN
In addition to the formation mechanism of H-Ti2AlN via the reaction of γ-TiAl with N
element according to Equ. (4), it is also possible to form H-Ti2AlN from α2-Ti3Al when both N and O are simultaneously present, which is discussed below. According to the calculated Ti-Al-O-N phase diagram at 900°C at a fixed ratio of xAl/xTi=1, H-Ti2AlN could exist in equilibrium with α-Al2O3 at a specific ratio of p O2 p N 2 [50]. Considering that only little residues of α2-Ti3Al were observed at the sub-interface which were severely un-matched with 16
ACCEPTED MANUSCRIPT bulk mass of Z-phase, and the phases in the vicinity of H-Ti2AlN are TiO2, Al2O3, AlNb2, Laves phase and the remaining α2-Ti3Al, one would expect that H-Ti2AlN could come from the decomposition of α2-Ti3Al in the concurrent presence of both N and O elements, i.e.,
Ti 3Al + N + 2O → Ti 2 AlN + TiO 2 .
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(8)
Since no TiO2 adjacent to Z-phase or original γ-TiAl/α2-Ti3Al lamellae was observed in Fig. 7(a), it could be assumed that the outer equiaxed TiO2 grains had formed through the
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following two paths: (a) As reaction (8) shows, Ti would diffuse outwards in Ti3Al to react
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with O, and (b) Ti2AlN would be oxidized to form TiO2 and Al2O3 as expressed by the following equation,
2Ti 2 AlN + 11O → 4TiO 2 + Al2O3 + 2 N .
(9)
Considering that oxygen partial pressure would increase at the nitride layer as the
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interface moved inwards [22,49], the N released from the above oxidation reaction of Ti2AlN would further promote the nitridation of α2-Ti3Al as expressed by Equ. (8).
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4.6 Oxidation mechanism of TiAlNbCr alloy
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The oxidation of α2-Ti3Al and γ-TiAl intermetallic alloys was reported to include three stages: absorption of oxygen, selective growth of alumina and growth of Ti oxides [58]. Although Al2O3 grew preferentially in the initial stage, TiO2 grains with a good deal of oxygen vacancies overgrew with respect to the Al2O3 grains because of the faster growth rate as seen from the outermost coarse TiO2 layer on the surface in Fig. 4(a). Scarcely any niobium was observed in this layer, which means that TiO2 formed by the outward diffusion of Ti, as also observed in [59]. The Al2O3-rich layer (dark area in Fig. 4(a)) beneath TiO2 17
ACCEPTED MANUSCRIPT (which was not dense enough to prevent the diffusion of O and N) was formed when the relative activity of Al rose after Ti was consumed. The formation of outer Al2O3 and TiO2 film in conjunction with inward diffusion of N and O during oxidation of TiAlNbCr at 900°C
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could be summarized in Fig. 12. Stage “1” represented the formation of the outmost Al2O3 and TiO2 film in conjunction with inward diffusion of N and O. In stage “2”, H-Ti2AlN/α2-Ti3Al lamellae came into being underneath the oxide film from the original
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γ-TiAl/α2-Ti3Al lamellae owing to the decrease of ratio p O2 p N 2 . Then in stage “3” cubic
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Z-phase formed at the expense of H-Ti2AlN/α2-Ti3Al lamellae. In the subsequent oxidation of stage “4”, Z-phase decomposed into Al2O3 and α2-Ti3Al which further reacted with both O and N to form H-Ti2AlN. Finally, in stage “5” H-Ti2AlN decomposed into TiO2 and Al2O3,
5. Conclusions
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leading to the formation of a TiO2-rich layer as the scale/substrate interface moved inward.
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(1) The microstructure of a newly-developed TiAlNbCr alloy consisted of characteristic
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γ-TiAl/α2-Ti3Al lamellae with B2-phase and γ-phase present at the grain boundaries. The alloy exhibited a superior oxidation resistance at 700°C, and its high-temperature oxidation kinetics at 800°C and 900°C followed well the parabolic law. (2) Coarse columnar rutile TiO2 formed on the outermost surface after oxidation at high temperatures. γ-TiAl in the γ/α2 lamellae in the vicinity of oxide scale/substrate interface was observed to transform into hexagonal H-Ti2AlN, while α2-Ti3Al remained unchanged. The original γ/α2 lamellae near the scale/substrate interface were thus substituted by the 18
ACCEPTED MANUSCRIPT formation of H-Ti2AlN/α2-Ti3Al lamellae, with an orientation relationship identified to be (0001)α2 //(0001)Ti2AlN , (10 1 0)α 2 //(10 1 0)Ti 2AlN and [1210]α 2 //[1210]Ti 2AlN . (3) The oxide scale exhibited a highly-complex multi-layered structure consisting of TiO2, H-Ti2AlN+Al2O3+α2-Ti3Al,
Al2O3+TiO2,
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Al2O3-rich,
H-Ti2AlN/α2-Ti3Al
Z-Ti5Al3O2+AlNb2+Laves-(Ti,Nb)Cr2,
lamellae
from
the
outside to inside after high-temperature oxidation. The metastable Z-Ti5Al3O2 phase at
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the scale/substrate interface came from the oxidation of H-Ti2AlN/α2-Ti3Al lamellae and
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then it was decomposed to Ti3Al and Al2O3 as the scale/substrate interface moved inwardly during oxidation. Then Ti3Al reacted with oxygen and nitrogen to become Ti2AlN, which was finally oxidized to form TiO2 and Al2O3.
(4) Nb was doped in the rutile to suppress the diffusion of oxygen, and this doping effect
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acted mainly in the TiO2 of the TiO2+Al2O3 compound layer. Nb along with Cr and Mo transformed to AlNb2 and Laves phase at the scale/substrate interface, which blocked the formation of continuous Z-phase and reduced the direct exposure of Z-phase and
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adjacent H-Ti2AlN/α2-Ti3Al lamellae to oxygen and nitrogen.
Acknowledgments
The authors are grateful for the financial support provided by the National Natural Science Foundation of China (NSFC) (Grant Nos. 51305304 and U1302275) and the Natural Sciences and Engineering Research Council of Canada (NSERC) in the form of international research collaboration, Fundamental Research Funds for the Central Universities, and Major Science and Technology Project “High-end CNC Machine Tools and Basic Manufacturing Equipment” (2013ZX04011061). The authors also thank Prof. G.J. Cao of Harbin University 19
ACCEPTED MANUSCRIPT of Science and Technology, for his assistance in the FIB and TEM/HREM experiments.
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Table Captions:
Table 1 Local chemical compositions (in at. %) at various points indicated in Fig. 4(a) for the TiAlNbCr alloy. after oxidation at 900°C for 100 h.
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Table 2 Crystallographic lattice parameters of Ti2AlN, Ti3Al and TiAl. Table 3 Fitting results of free energy curves of Ti2AlN, Ti3Al and TiAl.
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Figures Captions:
Figure 1 Microstructure of as-cast TiAlNbCr alloy: (a) SEM-BSE micrograph, (b) TEM-BF
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image along with the SAD patterns of points A (α2-Ti3Al) and B (γ-TiAl). Figure 2 Oxidation kinetics of TiAlNbCr alloy at 700, 800 and 900°C up to 100 h. Figure 3 XRD spectra of oxide scales in TiAlNbCr alloy after isothermal oxidation for 100 h at (a) 700°C, (b) 800°C and (c) 900°C.
Figure 4 Cross-sectional morphology of TiAlNbCr alloy after 100 h of oxidation at 900°C: (a) SEM micrograph and (b) EPMA elemental mapping across the oxide scale. Figure 5 A TEM/STEM sample of TiAlNbCr alloy oxidation at 900°C for 100 h, prepared via focused-ion-beam (FIB) technique. Figure 6 TiAlNbCr alloy after oxidation at 900°C for 100 h: (a) TEM BF image at the 23
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magnification image of zone “A” in (a), and (l) HRTEM image of point 10 in (k). Figure 8 TiAlNbCr alloy after oxidation at 900°C for 100 h: (a) HAADF image of a sub-interface area, (b) and (c) higher magnification images of zones “A” and “B” in (a), and (d) and (e) SAD patterns of points “ Ι ” and “ ΙΙ ”.
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Figure 9 TiAlNbCr alloy after oxidation at 900°C for 100 h: (a) BF image of lamellae close to scale/substrate interface, and (b)-(f) SAD patterns of points A-E, respectively. Primitive unit cell structure of (a) γ-TiAl, (b) α2-Ti3Al and (c) Ti2AlN.
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Figure 10
Figure 11 Gibbs free energy curves of γ-TiAl, α2-Ti3Al, and H-Ti2AlN as a function of temperature before and after fitting.
Figure 12 Schematic illustration for the formation of multi-layered oxide scale during
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oxidation of TiAlNbCr alloy at 900°C.
24
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Ti
Al
Nb
Cr
Mo
B
Y
N
O
1
45.42
0.97
0.03
0.05
0
0
49.23
2
35.29
0.11
0
0.04
0
0
5.24
59.32
3
3.19
35.63
0
0.15
0
0.40
1.81
40.27
0
0.18
0.01
0.08
5
15.62
18.13
1.24
0.48
0
0
6
26.42
3.53
2.12
0.53
0
0
7
30.55
0.98
1.43
0.47
0
0
8
35.61
20.53
5.21
1.89
0.43
0
9
28.40
14.92
5.08
1.83
0.40
0
10
47.65
31.99
6.86
1.39
0.44
0
11
47.70
32.89
5.70
1.30
0.46
12
53.52
28.46
3.90
1.23
0.18
13
51.18
26.79
3.69
1.11
0.22
14
54.28
16.83
3.42
15
47.55
26.69
5.90
16
49.12
21.16
5.16
0.44 1.00 1.37 3.77 4.79 19.25 15.61 6.50 7.25 7.46 8.04 8.10 7.82 7.31
60.19
4
0 0 0.01 0.02 0.01 0 0 0.01 0 0.03 0.02 0.05 0 0 0 0
4.31
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oxidation at 900°C for 100 h.
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3.21
2.11
0
3.17
1.69
0
7.07
3.38
0
56.62 63.16 63.65 61.78 17.08 33.77 5.15 4.67 5.19 8.98 12.01 7.20 6.80
Table 2 Crystallographic lattice parameters of Ti2AlN, Ti3Al and TiAl. Crystal
group
system
#52641
#99779
TiAl
#99780
P 63/mmc
Hexagonal
P 63/mmc
Hexagonal
P 4/mmm
Tetragonal
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Ti3Al
Site of atoms
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Space
EP
Ti2AlN
ICSD number
x
y
z
Ti
4f
0.33333
0.66667
0.5860
Al
2c
0.33333
0.66667
0.25
N
2a
0
0
0
Ti
6h
0.16670
0.33340
0.25
Al
2d
0.33333
0.66667
0.75
Ti
1a
0
0
0
Al
1d
0.5
0.5
0.5
Table 3 Fitting results of free energy curves of Ti2AlN, Ti3Al and TiAl.
Phases
Adjusted
H0
a
b
c
I
Ti2AlN
6.03866
-0.02759
-9.96743E-6
45.03363
0.15504
0.99999
Ti3Al
2.53077
-0.03492
-5.17107E-6
49.26238
0.19305
1
TiAl
0.51999
-0.00457
-1.04607E-6
8.2674
0.02637
0.99999
25
R-squared
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Figure 1 Microstructure of as-cast TiAlNbCr alloy: (a) SEM-BSE micrograph, (b) TEM-BF image along with the SAD patterns of points A (α2-Ti3Al) and B (γ-TiAl).
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Figure 2 Oxidation kinetics of TiAlNbCr alloy at 700, 800 and 900°C up to 100 h.
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Figure 3 XRD spectra of oxide scales in TiAlNbCr alloy after isothermal oxidation for 100 h at (a) 700°C, (b) 800°C and (c) 900°C.
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Figure 4 Cross-sectional morphology of TiAlNbCr alloy after 100 h of oxidation at 900°C: (a) SEM micrograph and (b) EPMA elemental mapping across the oxide scale.
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Figure 5 A TEM/STEM sample of TiAlNbCr alloy oxidation at 900°C for 100 h, prepared via focused-ion-beam (FIB) technique.
Figure 6 TiAlNbCr alloy after oxidation at 900°C for 100 h: (a) TEM BF image at the scale/substrate interface, (b) the corresponding HAADF image, and (c)-(k) STEM-EDS elemental mappings of Ti, Al, O, N, Nb, Cr, Mo, B , Y, respectively.
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Figure 7 TiAlNbCr alloy after oxidation at 900°C for 100 h: (a) BF image near the interface of scale and lamellar substrate, (b)-(j) SAD patterns of points 1-9, (k) higher magnification image of zone “A” in (a), and (l) HRTEM image of point 10 in (k).
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Figure 8 TiAlNbCr alloy after oxidation at 900°C for 100 h: (a) HAADF image of a sub-interface area, (b) and (c) higher magnification images of zones “A” and “B” in (a), and (d) and (e) SAD patterns of points “ Ι ” and “ ΙΙ ”.
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Figure 9 TiAlNbCr alloy after oxidation at 900°C for 100 h: (a) BF image of lamellae close to scale/substrate interface, and (b)-(f) SAD patterns of points A-E, respectively.
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Figure 10 Primitive unit cell structure of (a) γ-TiAl, (b) α2-Ti3Al and (c) Ti2AlN.
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Figure 11 Gibbs free energy curves of γ-TiAl, α2-Ti3Al, and H-Ti2AlN as a function of temperature before and after fitting.
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Figure 12. Schematic illustration for the formation of multi-layered oxide scale during oxidation of TiAlNbCr alloy at 900°C.
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S1359-6454(18)30114-9
DOI:
10.1016/j.actamat.2018.02.013
Reference:
AM 14369
To appear in:
Acta Materialia
Received Date: 26 November 2017 Revised Date:
28 January 2018
Accepted Date: 5 February 2018
Please cite this article as: S.J. Qu, S.Q. Tang, A.H. Feng, C. Feng, J. Shen, D.L. Chen, Microstructural evolution and high-temperature oxidation mechanisms of a titanium aluminide based alloy, Acta Materialia (2018), doi: 10.1016/j.actamat.2018.02.013. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
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ACCEPTED MANUSCRIPT Microstructural evolution and high-temperature oxidation mechanisms of a titanium aluminide based alloy
S.J. Qu1*, S.Q. Tang1, A.H. Feng1, C. Feng1, J. Shen1, D.L. Chen2∗ School of Materials Science and Engineering, Tongji University, Shanghai 201804, China 2
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1
Department of Mechanical and Industrial Engineering, Ryerson University, Toronto,
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Ontario M5B 2K3, Canada
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Abstract: Oxidation resistance of titanium aluminide (TiAl) based alloys is a fundamental aspect for the high-temperature structural applications such as in the advanced hypersonic aircraft engines and gas turbines. The aim of this study was to identify oxidation kinetics and mechanisms through detailed microstructural characterization of a newly-developed
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Ti-44Al-4Nb-1.5Cr-0.5Mo-0.1B-0.1Y alloy via focused ion beam (FIB), transmission electron microscopy (TEM), X-ray diffraction (XRD), electron probe microanalysis (EPMA),
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scanning transmission electron microscopy (STEM), along with density functional theory (DFT) calculations. The alloy consisting mainly of γ-TiAl/α2-Ti3Al lamellar structure
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exhibited a superior oxidation resistance at 700°C, and followed parabolic oxidation kinetics at 800°C and 900°C. The observed multi-layered scale structure consisted of TiO2, Al2O3-rich, Al2O3+TiO2,
H-Ti2AlN+Al2O3+α2-Ti3Al,
Z-Ti5Al3O2+AlNb2+Laves-(Ti,Nb)Cr2,
and
H-Ti2AlN/α2-Ti3Al lamellae from the outside to inside after high-temperature oxidation. The γ-TiAl/α2-Ti3Al lamellae near the scale/substrate interface were first transformed into
∗
Corresponding author. Tel.: +86-21-6958-1009; E-mail address: [email protected] (S.J. Qu); and Tel: (416) 979-5000 ext. 6487; Fax: (416) 979-5265; Email: [email protected] (D.L. Chen). 1
ACCEPTED MANUSCRIPT H-Ti2AlN/α2-Ti3Al lamellae, with orientation relationships identified as (0001)α2 //(0001)Ti2AlN , (10 1 0)α 2 //(10 1 0)Ti 2AlN and [12 10]α 2 //[1210]Ti 2AlN . The H-Ti2AlN/α2-Ti3Al lamellae were
then transformed into a metastable Z-Ti5Al3O2 phase at the scale/substrate interface. The
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Z-phase was decomposed to Ti3Al and Al2O3 as the scale/substrate interface moved inwardly. Ti3Al reacted further with oxygen and nitrogen to form Ti2AlN, which was finally oxidized to form TiO2 and α-Al2O3. A Nb-rich layer was present beneath the scale along with the
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formation of AlNb2 and Laves phase, and the doping effect of Nb to suppress the diffusion of
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oxygen occurred mainly in the TiO2+Al2O3 compound layer. The results obtained in this study would pave the way for the development of advanced oxidation-resistant TiAl-based materials for high-temperature applications.
1. Introduction
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Keywords: TiAl-based alloys; oxidation kinetics; oxidation mechanisms; EPMA; TEM.
As a relatively new class of promising materials for high-temperature structural
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applications such as in the advanced hypersonic aircraft engines and gas turbines,
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intermetallic titanium aluminide (TiAl) based alloys have recently drawn considerable attention due to their attractive characteristics, including lightweighting (with a low density of ρ≈3.9-4.2 g/cm3), high specific strength, excellent modulus retention, and superior creep resistance [1-6]. This is indeed motivated by the successful application of a TiAl-based alloy in General Electric’s high-thrust GEnex jet engine for powering Boeing 747-8 and 787 Dreamliner [2,3,7,8]. TiAl-based alloys in polycrystalline forms have been increasingly utilized to replace nickel-based superalloys (with a density of ρ≈8.5 g/cm3) in the temperature 2
ACCEPTED MANUSCRIPT range of 650-750°C, with the benefit of ~50% weight reduction [2,3,9,10]. However, the widespread commercial applications of these alloys are still limited due to the low room-temperature ductility and some challenges in manufacturing [1,11], and especially
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unsatisfactory oxidation resistance above ~750°C [12,13]. Because of a similar activity of Ti and Al in the severe environment of turbine engines, a protective Al2O3 layer could not be properly formed on the surface of TiAl-based alloys after
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high-temperature oxidation [14]. Generally, nonprotective multi-layered oxide scales
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composed of TiO2/Al2O3/TiO2+Al2O3 occurred after oxidation of TiAl, since γ-TiAl cannot be directly oxidized to TiO2 and Al2O3 [15-17]. During oxidation titanium nitrides (TiN and Ti2AlN) along with a metastable cubic Z-phase (Ti5Al3O2) (or X-phase) have been observed at the interface of oxide scale and substrate via TEM [18-20]. Copland et al. [21] reported
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that in binary alloys containing γ-TiAl and α2-Ti3Al phases, metastable Z-phase first occurred from the oxidation of γ-TiAl and eventually decomposed to α2-Ti3Al (rich in oxygen) and Al2O3 when oxidation experiments were conducted at 1000°C in a high-purity oxygen
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environment. Dettenwanger et al. [22] observed an Al-depletion layer after oxidation of a
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binary Ti-50Al alloy at 900°C for 1 h in air, where the Z-phase was also the first phase formed. In the subsequent oxidation, α2-Ti3Al will form at the interface of Z-phase and γ-TiAl. The Z-phase with a higher Vickers hardness may induce cracks in the subsurface zone and Al2O3 layer [21,22]. In addition, the Z-phase could destroy the protective Al2O3 layer after its decomposition to α2-Ti3Al and Al2O3 [21]. Titanium nitrides were observed only at the substrate/scale interface where the ratio of p O2 p N 2 is low, because the formation of the oxides leads to a reduction of p O 2 [15]. Nitrides would be oxidized to TiO2 after long 3
ACCEPTED MANUSCRIPT exposure and the released nitrogen can result in the formation of a new nitride layer. This repeated process thus impedes the formation of a continuous Al2O3 layer [22]. The orientation relationship between the nitrides in the nitride layer was observed to be (111)TiN || (0001)Ti AlN 2
and [1 11]TiN || [1120]Ti AlN . High-density stacking faults in the TiN phase and Al diffusion from
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2
the substrate to the outer layer lead to the formation of the Ti2AlN phase at the bottom of the TiN phase [23]. However, it is unclear how the Z-phase evolves into nitrides and if the
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lamellar structure below the Z-phase would change in the oxidation process.
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To hinder the diffusion of O and N from air to substrate, several measures were taken [24], such as coating technologies [25-27] and surface treatment [28-31]. Adding moderate ternary elements, such as Nb, W, Si, Mo, Cr and Y in TiAl-based alloys, improves not only the mechanical properties but also the oxidation resistance at high temperatures [32-37].
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Dopant elements with a valence higher than Ti4+, such as Nb5+, would cause a reduction of O vacancy concentration in the TiO2 lattice in order to maintain electrical neutrality. Thus the growth of TiO2 is blocked at a lower oxygen diffusion rate [38]. A continuous sub-scale
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Ti(Cr,Al)2 Laves phase could form if a certain amount (7~10 at.%) of Cr was added to
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TiAl-based alloys [39,40]. The addition of element Y results in a reduction of weight gain and an improvement of adhesion between oxide scale and substrate [34,41]. Although the role of the added elements was discussed, their influence on the microstructural evolution of Z-phase,
α2-Ti3Al and substrate was rarely identified. Besides, it is unknown which phase in the γ-TiAl/α2-Ti3Al lamellar structure would change during high-temperature oxidation. The purpose of this study was, therefore, to identify the microstructural evolution from
γ-TiAl/α2-Ti3Al to Al2O3 (corundum) and TiO2 (rutile) and fundamental oxidation 4
ACCEPTED MANUSCRIPT mechanisms via detailed characterization of the whole oxide scale after oxidation of a newly-developed TiAlNbCr alloy at 900°C for 100 h via SEM, STEM, EPMA, FIB and TEM,
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along with relevant thermodynamic calculations involving density functional theory (DFT).
2.1. Specimen preparation material
with
a
nominal
chemical
composition
of
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The
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2. Experimental procedures
Ti-44Al-4Nb-1.5Cr-0.5Mo-0.1B-0.1Y (at.%) (hereafter referred to as TiAlNbCr alloy) was prepared by double vacuum consumable arc melting technique. For oxidation tests, the specimens with approximate dimensions of 10×10×1 mm3 were cut from the alloy ingots by
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wire EDM. The surfaces were ground using SiC sandpapers up to grit No. 1200, and then cleaned ultrasonically in acetone for 15 min. Finally, the initial size and weight of each
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specimen were measured for calculating weight gain later.
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2.2. Oxidation tests
Isothermal oxidation tests were carried out at 700, 800 and 900°C in a muffle furnace in static laboratory air for 100 h. A hole on the door with a diameter of 10 mm was drilled to ensure sufficient oxygen in the furnace chamber. After a specified temperature has reached, multiple specimens were put in the furnace and then individual specimens were taken out after 1, 3, 6, 12, 24, 36, 50, 62, 74, 86 and 100 h, respectively. Their weights were measured by an electronic balance with a resolution of 0.1 mg to calculate the weight gain ∆M. 5
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2.3. Microstructural characterization Original microstructures of the test material were examined by field emission
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environmental scanning electron microscope (FE-ESEM, FEI Quanta 250 FEG) equipped with energy-dispersive X-ray spectroscopy (EDS, Apollo SDD10), and transmission electron microscopy (TEM, FEI Talos 200X). X-ray diffraction (XRD, Rigaku D/Max-2550) with Cu
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Kα radiation (λ=1.5418Å) was used to identify the phases in the oxide scale at 50 kV and 200
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mA with a diffraction angle (2θ) from 10° to 100° with a step size of 0.02° and 1 s in each step. Elemental distribution in different regions on the cross-section of the oxide scale was characterized by using the electron probe microanalyzer (EPMA, Shimadzu 1720) with a resolution of 1 µm and secondary-electron image resolution of 6 nm using a beam current of
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10 nA, and scanning transmission electron microscopy (STEM, FEI Talos 200X) with EDS. In the EPMA the standard samples used for calibrating the elemental distributions of titanium, aluminum, oxygen, nitrogen, niobium, chromium, molybdenum, boron and yttrium were pure
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titanium, Al2O3, BN, pure niobium, pure chromium, pure molybdenum, pure boron and
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YP4O12, respectively. Phase constitution at the interface of scale and substrate was determined by focused ion beam (FIB, FEI NanoLab 600i) and TEM.
2.4 Computational details
Thermodynamic calculations were performed using Materials Studio DMol3 version 6.0.0 (Accelrys Software Inc.), a high-quality quantum mechanics program [42]. DMol3 uses density functional theory (DFT) with a numerical radial function [43] basis set to calculate 6
ACCEPTED MANUSCRIPT the electronic properties of molecules, surfaces and crystalline solid materials [44] from the first-principles. Geometrical optimizations and frequency calculations were carried out with the GGA-BLYP function in conjunction with a double numerical plus d-functions (DND)
3. Results
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3.1. Original microstructures of TiAlNbCr alloy
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basis set to calculate free energy (G) as a function of temperature.
Back-scattered electron (BSE) SEM micrograph and TEM bright field (BF) microstructure of as-cast TiAlNbCr alloy are shown in Fig. 1(a) and (b). The microstructure consisted of grains (or colonies) containing characteristic (γ-TiAl + α2-Ti3Al) lamellae, where
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the bright B2-phase (ordered β-Ti) and dark γ-phase were present at the grain boundaries (Fig. 1(a)). The representative γ+α2 lamellar structure could be better seen from the TEM image
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along with the selected area diffraction (SAD) patterns in Fig. 1(b).
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3.2. Isothermal oxidation kinetics
Curves of isothermal-oxidation kinetics of TiAlNbCr alloy at 700, 800 and 900°C are
shown in Fig. 2. The obtained mass gain of this alloy after 100 h at 700, 800 and 900°C was about 0.0033, 0.4482 and 2.3777 mg/cm2, respectively. Thus this alloy exhibited a superior oxidation resistance at 700°C, with respect to the situation at 800 and 900°C. To identify which law of oxidation kinetics is followed, the obtained experimental data was fitted using the following equation, 7
ACCEPTED MANUSCRIPT ∆M n = k n t ,
(1)
where ∆M represents weight gain per unit area (mg/cm2), n is an oxidation exponent (n=1, linear relationship; n=2, parabolic relationship), kn is a rate constant (mgn/cm2n h) and t is
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oxidation time (h). The fitting curves via Origin software were plotted in Fig. 2, and the oxidation exponent (n) at 800 and 900°C was obtained to be 1.83 and 1.80, respectively. Hence the oxidation at 800 and 900°C obeyed a parabolic growth kinetics (n≈2) from 0 to
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100 h. Under this circumstance the value of kn would reflect the oxidation rate. Compared
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with kn=2.44×10-3 mg2/cm4 h at 800°C, the oxidation rate constant at 900°C was much higher (kn = 4.99×10-2 mg2/cm4 h). It suggests that this alloy exhibited a poor oxidation resistance at 900°C. Therefore, the subsequent study will be mainly focused on the oxidation of this alloy
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at 900°C.
3.3. XRD results
X-ray diffraction patterns of oxide scales in TiAlNbCr alloy are shown in Fig. 3. It is
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seen that after oxidation at 700°C for 100 h, except the γ, α2 and B2 in the substrate, oxides
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could barely be detected, while the peaks of oxides (rutile TiO2 and corundum α-Al2O3) and nitrides (TiN and Ti2AlN) were obvious at 800°C. When the temperature increased to 900°C, the peaks of rutile and corundum were further elevated, implying an increasing extent of oxidation. Furthermore, while the peaks of Ti2AlN became stronger, the peaks of TiN disappeared at 900°C.
3.4. Cross-sectional observations 8
ACCEPTED MANUSCRIPT Fig. 4(a) presents the SEM-BSE micrograph of oxide scale after 100 h of oxidation at 900°C. According to the elemental distribution maps of the cross-section of the oxide scale detected by EPMA shown in Fig. 4(b), it is seen that the outermost grey layer I consisted of
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columnar particles, and the adjacent black layer II was enriched in O and Al. Al and Ti coexisted in a mixed grey and black layer III. Combined with the XRD patterns shown in Fig. 3, the scale structure formed was identified to be in the order of TiO2 layer (I)/Al2O3-rich
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layer (II)/mixed Al2O3 +TiO2 layer (III) from the outside to inside.
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To reveal the chemical composition quantitatively in various areas of the cross-sectional scales, EPMA point microanalyses were conducted, as indicated in Fig. 4(a). The obtained local chemical compositions of points 1-16 are summarized in Table 1. Nb was rich in the white layer (about 3~4 µm) beneath the interface of oxide scale and substrate as indicated by
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points 10 and 11 in Fig. 4(a) and Nb distribution map in Fig. 4(b). In the oxide scale, Nb was mainly distributed in the layer III of TiO2 + Al2O3 mixture (points 5-9), but absent in the outer TiO2 layer I (points 1 and 2) and mid Al2O3-rich layer II (points 3 and 4). According to the
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chemical compositions at points 8 and 9, a small amount of nitrides was presumably formed
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at the scale/substrate interface because a sharp increase of N content was also detected (Table 1), which will be confirmed via subsequent TEM examinations. Besides, N and O with a concentration of 5~9% permeated into the lamellar colonies in the vicinity of scale/substrate interface, as indicated by points 12 and 13 (Table 1), which might cause a potential reduction of ductility of the alloy. Cr was enriched along with Mo and Nb in the white areas marked by points 14, 15 and 16 at the grain boundaries close to the scale, which corresponded to the presence of B2 phase in the substrate. B2 phase could 9
ACCEPTED MANUSCRIPT promote the formation of Al2O3, since the diffusion of Al was faster in B2 phase than in
γ-TiAl [45].
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3.5 TEM and STEM investigations The TEM/STEM sample of TiAlNbCr alloy after oxidation at 900°C for 100 h was carefully positioned at the scale/substrate interface, which was prepared via a
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focused-ion-beam (FIB) technique, as shown in Fig. 5. Fig. 6(a) and (b) show a TEM-BF
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image and a high-angle annular dark-field (HAADF) image at the scale/substrate interface, respectively. The lamellar structure at the bottom of Fig. 6(a) and (b) reflected the alloy substrate, and the equiaxed ultra-fine crystallites located above the lamellae were the microstructure of oxide scale. The elemental distribution in the entire area was very
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non-uniform as seen from Fig. 6(c)-(k), where the STEM-EDS elemental mapping of Ti, Al, O, N, Nb, Cr, Mo, B and Y was shown. The oxide scale mainly contained Ti, Al and O, and beneath the oxides there existed a distinct N-rich layer (Fig. 6(f)), which superimposed nicely
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with the location of Ti (Fig. 6(c)) and was in good agreement with the EPMA microanalysis
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results of points 8 and 9 in Fig. 4(a) and Table 1. Similarly, the positions of Al and O were also well overlapped, as seen from Fig. 6(d) and (e). The white particles present at the scale/substrate interface as seen from Fig. 6(b) were rich in Nb, Cr and Mo (Fig. 6(g)-(i)). More detailed TEM and high-resolution transmission electron microscopy (HRTEM) examinations of the interfacial area between the oxide scale and lamellar substrate were performed, as shown in Fig. 7. The oval area marked by “1” at the bottom-left of Fig. 7(a) had an obviously different morphology compared with the adjoining inner lamella and outer 10
ACCEPTED MANUSCRIPT ultra-fine oxide particles. The SAD pattern shown in Fig. 7(b) indicated that this oval block consisted of a simple cubic Z-phase (Ti5Al3O2), which exhibited a significantly higher Vickers hardness than base metal [22]. Another Z-phase was spotted in the area marked by
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“2”, which was far from area “1”. Areas “3” and “4” were identified as hexagonal H-phase (Ti2AlN) [46-48] through the SAD patterns (Fig. 7(d) and (e)). These discontinuous nitrides were present only at the scale/substrate interface, as also shown in Fig. 6(f). A characteristic
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microstructure of islands in area “5” consisted of alumina (α-Al2O3) formed besides the
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nitrides at the scale/substrate interface, corresponding well to Al and O mapping in Fig. 6(d) and (e). This suggests that the growth of α-Al2O3 oxides was interrupted by nitrides, making it difficult or impossible to form a dense protective layer, as also noted by Lu et al. [49]. The porous oxide layer outside the interface was mainly composed of a great deal of ultra-fine
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equiaxed rutile TiO2 such as areas “7” and “8” and a small amount of alumina (α-Al2O3, as indicated by area “6”). The area close to Z-phase marked by “9” was identified to be α2-Ti3Al and the other equiaxed α2 was only observed in area “10” in Fig. 7(k), which is a magnified
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image of zone “A” boxed in Fig. 7(a).
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Furthermore, the white particles at the scale/substrate interface shown in Fig. 8(a) were rich in Nb, Cr and Mo according to the STEM-EDS results in Fig. 6(g)-(i). Fig. 8(b) and (c) show higher magnification HAADF images of zones “A” and “B” in Fig. 8(a), where point “ Ι ” in Fig. 8(b) was enriched in Nb, i.e., the phase of AlNb2 as revealed by the SAD results (Fig. 8(d)), and point “ ΙΙ ” rich in Nb and Cr in Fig. 8(c) was identified as Laves phase (Nb, Ti)Cr2 (Fig. 8(e)). While the lamellar structure in the neighborhood of scale/substrate interface in Fig. 9(a) 11
ACCEPTED MANUSCRIPT looked seemingly like that in Fig. 1(b), the SAD results revealed that the outermost layer marked “A” of this specific lamella was neither α2-Ti3Al nor γ-TiAl, but the H-phase (Ti2AlN). SAD pattern from point “B” (i.e., the interface of two lamellae) showed that the
between
these
two
phases
were
(0001)α2 //(0001)Ti2AlN ,
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lamella adjacent to “A” still remained to be α2-Ti3Al, and the orientation relationships (10 1 0) α2 //(10 1 0) Ti2AlN
and
[1210]α 2 //[1210]Ti 2AlN . The inner layers “C” and “D” corresponded to H-Ti2AlN and α2-Ti3Al,
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respectively, and point “E” at some distance from the scale/substrate interface still had the
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same SAD pattern as point “B”. As a result, there was no γ-TiAl detected in the visible lamellar colony near the scale/substrate interface, while the α2-Ti3Al lamella remained unchanged after oxidation. This suggests that the original γ-TiAl/α2-Ti3Al lamellar structure adjacent to the scale was replaced by the H-Ti2AlN/α2-Ti3Al lamellar structure after oxidation
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4. Discussion
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at 900°C for 100 h.
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4.1 Multi-layered microstructural characteristics of oxide scale After isothermal oxidation in laboratory air at 900°C for 100 h, the oxide scale formed on the surface of TiAlNbCr alloy with a thickness of ~13 µm, which consisted of multiple layers of different oxides, nitride and intermetallic compounds. Based on the above examinations, the multi-layers from the outside to the inside were TiO2, Al2O3-rich, Al2O3+TiO2,
H-Ti2AlN+Al2O3+α2-Ti3Al,
Z-Ti5Al3O2+AlNb2+Laves-(Ti,Nb)Cr2,
H-Ti2AlN/α2-Ti3Al lamellae, and then alloy substrate (i.e., original γ-TiAl/α2-Ti3Al lamellar 12
ACCEPTED MANUSCRIPT structure). TiO2 and Al2O3 were the main oxidation products because they were thermodynamically stable in the Ti-Al-O-N system when oxygen was sufficiently offered [50]. TiO2 was a fast-growing oxide that did not provide long-term oxidation protection [51],
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consequently it occupied the largest scale volume. Discontinuous Z-phase, AlNb2 and Laves phase were present at the scale/substrate interface beneath the Ti2AlN-rich layer. The
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formation and decomposition of these phases will be discussed later.
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4.2 Oxidation kinetics
TiAlNbCr alloy exhibited a superior oxidation resistance at 700°C, as seen from Fig. 2, but weight gains became increasingly obvious at 800 and 900°C which basically obeyed a parabolic law. This suggests that the growth of oxide scale was controlled by diffusion
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processes because the parabolic law could be derived from Fick's first law of diffusion. During the parabolic oxidation of TiAl-based alloys, most of the scale formed by inward diffusion of oxygen [52]. TiO2 was the non-metal deficient n-type oxide with many oxygen
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vacancies and Ti4+ interstitial cations and thus became the main diffusion path [38]. The
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doping of Nb5+ was expected to decrease the oxygen vacancy concentration in order to maintain electroneutrality in TiO2 [53]. The relationship between diffusion coefficient D and the vacancy concentration Cv could be expressed by [54],
D = ACv F ,
(2)
where A is a crystal lattice-related parameter and F is the successful jump frequency of the atom. It is clear that oxygen diffusion would be suppressed when oxygen vacancies decreased. Jiang et al. [55] also reported that the oxidation resistance of TiAl-Nb increased with 13
ACCEPTED MANUSCRIPT increasing concentration of Nb from 0 to 20%. Therefore, the addition of Nb in the present alloy could decrease oxidation rate. That is, the observed white Nb-rich layer beneath the scale in Fig. 8(a) could be considered as a Nb-rich reservoir to block the inward diffusion of
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oxygen to a certain extent. However, further studies in this aspect are needed.
4.3 Formation of H-Ti2AlN/α2-Ti3Al lamellae
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It is of particular interest to observe that the H-Ti2AlN/α2-Ti3Al lamellar structure formed
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under the layer of Z-phase and Nb, Cr-rich phases (AlNb2 and Laves) to replace the original
γ-TiAl/α2-Ti3Al lamellae (Fig. 9). To understand this phenomenon, module DMol3 in the software of Material Studio 6.0 was used to assess thermodynamic possibility of phase transformation from γ-TiAl or α2-Ti3Al to H-Ti2AlN. Primitive crystal structures of γ-TiAl,
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α2-Ti3Al and H-Ti2AlN which are needed in the DFT calculations are shown in Fig. 10, with their lattice constants and atomic coordinates from ICSD database listed in Table 2. The dots in Fig. 11 represent the change of free energy as a function of temperature from 0-1000 K for
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γ-TiAl, α2-Ti3Al and H-Ti2AlN calculated via DMol3, and the corresponding curves were
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fitted according to the following equation, G = H 0 + aT ln T + bT 2 + c T + IT ,
(3)
where a, b, c, and I are all fitting parameters. The results obtained via the fitting are shown in Table 3 with an adjusted R-squared very close to 1. The Gibbs free energy change based on the following reactions,
4TiAl + 2 N → 2Ti 2AlN + 2Al ,
(4)
2Ti 3Al + 2 N → 2Ti 2 AlN + 2Ti ,
(5)
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ACCEPTED MANUSCRIPT at 900°C are obtained to be -77.6 kJ/mol and 26.2 kJ/mol, respectively. This indicates that the transformation of γ-TiAl to H-Ti2AlN was thermodynamically possible, while the transformation of α2-Ti3Al to H-Ti2AlN was impossible. Therefore, there was a strong
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tendency of transformation from γ-TiAl/α2-Ti3Al lamellae to H-Ti2AlN/α2-Ti3Al lamellae due to the presence of thermodynamic driving force at 900°C. Indeed, such a transformation
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would be possible when the temperature was above ~420 K, as seen from Fig. 11.
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4.4 Formation and decomposition of Z-phase
The metastable Z-phase with a chemical formula of Ti5Al3O2 was observed at the scale frontier, i.e., at the interface of scale and H-Ti2AlN/α2-Ti3Al lamellae (Fig. 7). Dettenwanger et al. [22] and Copland et al. [21] reported that for binary Ti-50Al alloy and Ti-52Al
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single-phase γ-TiAl alloy, the Z-phase formed in the depletion layer stayed continuous after oxidation at 900°C for 100 h and at 1000°C for 50 h, respectively. According to Lu et al. [49] and Copland et al. [21], the Z-phase came from the oxygen absorption of substrate alloy, but
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our TEM observations showed that the Z-phase was present discontinuously, which was
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surrounded by H-Ti2AlN/α2-Ti3Al lamellae, AlNb2, α-Al2O3 and H-Ti2AlN as shown in Fig. 7(a) and Fig. 8(a). Then the only possible source of Z-phase would be related to the H-Ti2AlN/α2-Ti3Al lamellae rather than the substrate γ-TiAl/α2-Ti3Al lamellae, and the formation process could be expressed as follows (where the chemical formula of H-Ti2AlN/α2-Ti3Al
lamellae
could
be
re-written
as
[50Ti-44Al-4Nb-1.5Cr-0.5Mo-0.1B-0.1Y)]xN1-x in our material), [50Ti -44Al-4Nb-1.5Cr -0.5Mo-0.1B-0.1Y )]x N (1− x ) + 20 xO → 10 xTi5 Al3O2 + 14 xAl + 4 xNb + 1.5xCr+0.5xMo+0.1xB+0.1xY + (1-x)N . (6) 15
ACCEPTED MANUSCRIPT This suggests that there appeared an enrichment of Al, Nb and N after the formation of Z-phase. Then AlNb2 and Laves phase of (Ti, Nb)Cr2 would form close to the Z-phase as shown in Fig. 8(a) and these two phases would impede continuous Z-phase formation in the
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further oxidation [49]. Laves phase (Ti, Nb)Cr2 had a good resistance to oxidation [56], and thus exposure of Z-phase and H-Ti2AlN/α2-Ti3Al lamellae to the oxygen could be alleviated to some extent because of the presence of AlNb2 and Laves phase at the sub-interface. It is
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expected that the oxidation resistance of TiAl-based alloys could be greatly improved if the
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added Nb and Cr was sufficient to form a continuous layer consisting of AlNb2 and Laves phases. The N atom released according to Equ. (6) would continue to diffuse inwards to react with inner γ-TiAl to form H-Ti2AlN based on Equ. (4), further corroborating the formation of H-Ti2AlN/α2-Ti3Al lamellae, as discussed above. It should be noted that the Z-phase could
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exist in equilibrium with γ-TiAl, α2-Ti3Al and α-A12O3 at specific oxygen concentrations [57]; it could also decompose in the following oxidation process [21,49], (7)
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3Ti 5 Al3O 2 → 2Al2O3 + 5Ti 3Al .
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4.5 Formation and decomposition of H-Ti2AlN
In addition to the formation mechanism of H-Ti2AlN via the reaction of γ-TiAl with N
element according to Equ. (4), it is also possible to form H-Ti2AlN from α2-Ti3Al when both N and O are simultaneously present, which is discussed below. According to the calculated Ti-Al-O-N phase diagram at 900°C at a fixed ratio of xAl/xTi=1, H-Ti2AlN could exist in equilibrium with α-Al2O3 at a specific ratio of p O2 p N 2 [50]. Considering that only little residues of α2-Ti3Al were observed at the sub-interface which were severely un-matched with 16
ACCEPTED MANUSCRIPT bulk mass of Z-phase, and the phases in the vicinity of H-Ti2AlN are TiO2, Al2O3, AlNb2, Laves phase and the remaining α2-Ti3Al, one would expect that H-Ti2AlN could come from the decomposition of α2-Ti3Al in the concurrent presence of both N and O elements, i.e.,
Ti 3Al + N + 2O → Ti 2 AlN + TiO 2 .
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(8)
Since no TiO2 adjacent to Z-phase or original γ-TiAl/α2-Ti3Al lamellae was observed in Fig. 7(a), it could be assumed that the outer equiaxed TiO2 grains had formed through the
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following two paths: (a) As reaction (8) shows, Ti would diffuse outwards in Ti3Al to react
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with O, and (b) Ti2AlN would be oxidized to form TiO2 and Al2O3 as expressed by the following equation,
2Ti 2 AlN + 11O → 4TiO 2 + Al2O3 + 2 N .
(9)
Considering that oxygen partial pressure would increase at the nitride layer as the
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interface moved inwards [22,49], the N released from the above oxidation reaction of Ti2AlN would further promote the nitridation of α2-Ti3Al as expressed by Equ. (8).
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4.6 Oxidation mechanism of TiAlNbCr alloy
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The oxidation of α2-Ti3Al and γ-TiAl intermetallic alloys was reported to include three stages: absorption of oxygen, selective growth of alumina and growth of Ti oxides [58]. Although Al2O3 grew preferentially in the initial stage, TiO2 grains with a good deal of oxygen vacancies overgrew with respect to the Al2O3 grains because of the faster growth rate as seen from the outermost coarse TiO2 layer on the surface in Fig. 4(a). Scarcely any niobium was observed in this layer, which means that TiO2 formed by the outward diffusion of Ti, as also observed in [59]. The Al2O3-rich layer (dark area in Fig. 4(a)) beneath TiO2 17
ACCEPTED MANUSCRIPT (which was not dense enough to prevent the diffusion of O and N) was formed when the relative activity of Al rose after Ti was consumed. The formation of outer Al2O3 and TiO2 film in conjunction with inward diffusion of N and O during oxidation of TiAlNbCr at 900°C
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could be summarized in Fig. 12. Stage “1” represented the formation of the outmost Al2O3 and TiO2 film in conjunction with inward diffusion of N and O. In stage “2”, H-Ti2AlN/α2-Ti3Al lamellae came into being underneath the oxide film from the original
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γ-TiAl/α2-Ti3Al lamellae owing to the decrease of ratio p O2 p N 2 . Then in stage “3” cubic
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Z-phase formed at the expense of H-Ti2AlN/α2-Ti3Al lamellae. In the subsequent oxidation of stage “4”, Z-phase decomposed into Al2O3 and α2-Ti3Al which further reacted with both O and N to form H-Ti2AlN. Finally, in stage “5” H-Ti2AlN decomposed into TiO2 and Al2O3,
5. Conclusions
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leading to the formation of a TiO2-rich layer as the scale/substrate interface moved inward.
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(1) The microstructure of a newly-developed TiAlNbCr alloy consisted of characteristic
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γ-TiAl/α2-Ti3Al lamellae with B2-phase and γ-phase present at the grain boundaries. The alloy exhibited a superior oxidation resistance at 700°C, and its high-temperature oxidation kinetics at 800°C and 900°C followed well the parabolic law. (2) Coarse columnar rutile TiO2 formed on the outermost surface after oxidation at high temperatures. γ-TiAl in the γ/α2 lamellae in the vicinity of oxide scale/substrate interface was observed to transform into hexagonal H-Ti2AlN, while α2-Ti3Al remained unchanged. The original γ/α2 lamellae near the scale/substrate interface were thus substituted by the 18
ACCEPTED MANUSCRIPT formation of H-Ti2AlN/α2-Ti3Al lamellae, with an orientation relationship identified to be (0001)α2 //(0001)Ti2AlN , (10 1 0)α 2 //(10 1 0)Ti 2AlN and [1210]α 2 //[1210]Ti 2AlN . (3) The oxide scale exhibited a highly-complex multi-layered structure consisting of TiO2, H-Ti2AlN+Al2O3+α2-Ti3Al,
Al2O3+TiO2,
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Al2O3-rich,
H-Ti2AlN/α2-Ti3Al
Z-Ti5Al3O2+AlNb2+Laves-(Ti,Nb)Cr2,
lamellae
from
the
outside to inside after high-temperature oxidation. The metastable Z-Ti5Al3O2 phase at
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the scale/substrate interface came from the oxidation of H-Ti2AlN/α2-Ti3Al lamellae and
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then it was decomposed to Ti3Al and Al2O3 as the scale/substrate interface moved inwardly during oxidation. Then Ti3Al reacted with oxygen and nitrogen to become Ti2AlN, which was finally oxidized to form TiO2 and Al2O3.
(4) Nb was doped in the rutile to suppress the diffusion of oxygen, and this doping effect
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acted mainly in the TiO2 of the TiO2+Al2O3 compound layer. Nb along with Cr and Mo transformed to AlNb2 and Laves phase at the scale/substrate interface, which blocked the formation of continuous Z-phase and reduced the direct exposure of Z-phase and
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adjacent H-Ti2AlN/α2-Ti3Al lamellae to oxygen and nitrogen.
Acknowledgments
The authors are grateful for the financial support provided by the National Natural Science Foundation of China (NSFC) (Grant Nos. 51305304 and U1302275) and the Natural Sciences and Engineering Research Council of Canada (NSERC) in the form of international research collaboration, Fundamental Research Funds for the Central Universities, and Major Science and Technology Project “High-end CNC Machine Tools and Basic Manufacturing Equipment” (2013ZX04011061). The authors also thank Prof. G.J. Cao of Harbin University 19
ACCEPTED MANUSCRIPT of Science and Technology, for his assistance in the FIB and TEM/HREM experiments.
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Table Captions:
Table 1 Local chemical compositions (in at. %) at various points indicated in Fig. 4(a) for the TiAlNbCr alloy. after oxidation at 900°C for 100 h.
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Table 2 Crystallographic lattice parameters of Ti2AlN, Ti3Al and TiAl. Table 3 Fitting results of free energy curves of Ti2AlN, Ti3Al and TiAl.
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Figures Captions:
Figure 1 Microstructure of as-cast TiAlNbCr alloy: (a) SEM-BSE micrograph, (b) TEM-BF
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image along with the SAD patterns of points A (α2-Ti3Al) and B (γ-TiAl). Figure 2 Oxidation kinetics of TiAlNbCr alloy at 700, 800 and 900°C up to 100 h. Figure 3 XRD spectra of oxide scales in TiAlNbCr alloy after isothermal oxidation for 100 h at (a) 700°C, (b) 800°C and (c) 900°C.
Figure 4 Cross-sectional morphology of TiAlNbCr alloy after 100 h of oxidation at 900°C: (a) SEM micrograph and (b) EPMA elemental mapping across the oxide scale. Figure 5 A TEM/STEM sample of TiAlNbCr alloy oxidation at 900°C for 100 h, prepared via focused-ion-beam (FIB) technique. Figure 6 TiAlNbCr alloy after oxidation at 900°C for 100 h: (a) TEM BF image at the 23
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magnification image of zone “A” in (a), and (l) HRTEM image of point 10 in (k). Figure 8 TiAlNbCr alloy after oxidation at 900°C for 100 h: (a) HAADF image of a sub-interface area, (b) and (c) higher magnification images of zones “A” and “B” in (a), and (d) and (e) SAD patterns of points “ Ι ” and “ ΙΙ ”.
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Figure 9 TiAlNbCr alloy after oxidation at 900°C for 100 h: (a) BF image of lamellae close to scale/substrate interface, and (b)-(f) SAD patterns of points A-E, respectively. Primitive unit cell structure of (a) γ-TiAl, (b) α2-Ti3Al and (c) Ti2AlN.
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Figure 10
Figure 11 Gibbs free energy curves of γ-TiAl, α2-Ti3Al, and H-Ti2AlN as a function of temperature before and after fitting.
Figure 12 Schematic illustration for the formation of multi-layered oxide scale during
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oxidation of TiAlNbCr alloy at 900°C.
24
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Ti
Al
Nb
Cr
Mo
B
Y
N
O
1
45.42
0.97
0.03
0.05
0
0
49.23
2
35.29
0.11
0
0.04
0
0
5.24
59.32
3
3.19
35.63
0
0.15
0
0.40
1.81
40.27
0
0.18
0.01
0.08
5
15.62
18.13
1.24
0.48
0
0
6
26.42
3.53
2.12
0.53
0
0
7
30.55
0.98
1.43
0.47
0
0
8
35.61
20.53
5.21
1.89
0.43
0
9
28.40
14.92
5.08
1.83
0.40
0
10
47.65
31.99
6.86
1.39
0.44
0
11
47.70
32.89
5.70
1.30
0.46
12
53.52
28.46
3.90
1.23
0.18
13
51.18
26.79
3.69
1.11
0.22
14
54.28
16.83
3.42
15
47.55
26.69
5.90
16
49.12
21.16
5.16
0.44 1.00 1.37 3.77 4.79 19.25 15.61 6.50 7.25 7.46 8.04 8.10 7.82 7.31
60.19
4
0 0 0.01 0.02 0.01 0 0 0.01 0 0.03 0.02 0.05 0 0 0 0
4.31
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oxidation at 900°C for 100 h.
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3.21
2.11
0
3.17
1.69
0
7.07
3.38
0
56.62 63.16 63.65 61.78 17.08 33.77 5.15 4.67 5.19 8.98 12.01 7.20 6.80
Table 2 Crystallographic lattice parameters of Ti2AlN, Ti3Al and TiAl. Crystal
group
system
#52641
#99779
TiAl
#99780
P 63/mmc
Hexagonal
P 63/mmc
Hexagonal
P 4/mmm
Tetragonal
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Ti3Al
Site of atoms
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Space
EP
Ti2AlN
ICSD number
x
y
z
Ti
4f
0.33333
0.66667
0.5860
Al
2c
0.33333
0.66667
0.25
N
2a
0
0
0
Ti
6h
0.16670
0.33340
0.25
Al
2d
0.33333
0.66667
0.75
Ti
1a
0
0
0
Al
1d
0.5
0.5
0.5
Table 3 Fitting results of free energy curves of Ti2AlN, Ti3Al and TiAl.
Phases
Adjusted
H0
a
b
c
I
Ti2AlN
6.03866
-0.02759
-9.96743E-6
45.03363
0.15504
0.99999
Ti3Al
2.53077
-0.03492
-5.17107E-6
49.26238
0.19305
1
TiAl
0.51999
-0.00457
-1.04607E-6
8.2674
0.02637
0.99999
25
R-squared
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Figure 1 Microstructure of as-cast TiAlNbCr alloy: (a) SEM-BSE micrograph, (b) TEM-BF image along with the SAD patterns of points A (α2-Ti3Al) and B (γ-TiAl).
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Figure 2 Oxidation kinetics of TiAlNbCr alloy at 700, 800 and 900°C up to 100 h.
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Figure 3 XRD spectra of oxide scales in TiAlNbCr alloy after isothermal oxidation for 100 h at (a) 700°C, (b) 800°C and (c) 900°C.
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Figure 4 Cross-sectional morphology of TiAlNbCr alloy after 100 h of oxidation at 900°C: (a) SEM micrograph and (b) EPMA elemental mapping across the oxide scale.
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Figure 5 A TEM/STEM sample of TiAlNbCr alloy oxidation at 900°C for 100 h, prepared via focused-ion-beam (FIB) technique.
Figure 6 TiAlNbCr alloy after oxidation at 900°C for 100 h: (a) TEM BF image at the scale/substrate interface, (b) the corresponding HAADF image, and (c)-(k) STEM-EDS elemental mappings of Ti, Al, O, N, Nb, Cr, Mo, B , Y, respectively.
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Figure 7 TiAlNbCr alloy after oxidation at 900°C for 100 h: (a) BF image near the interface of scale and lamellar substrate, (b)-(j) SAD patterns of points 1-9, (k) higher magnification image of zone “A” in (a), and (l) HRTEM image of point 10 in (k).
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Figure 8 TiAlNbCr alloy after oxidation at 900°C for 100 h: (a) HAADF image of a sub-interface area, (b) and (c) higher magnification images of zones “A” and “B” in (a), and (d) and (e) SAD patterns of points “ Ι ” and “ ΙΙ ”.
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Figure 9 TiAlNbCr alloy after oxidation at 900°C for 100 h: (a) BF image of lamellae close to scale/substrate interface, and (b)-(f) SAD patterns of points A-E, respectively.
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Figure 10 Primitive unit cell structure of (a) γ-TiAl, (b) α2-Ti3Al and (c) Ti2AlN.
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Figure 11 Gibbs free energy curves of γ-TiAl, α2-Ti3Al, and H-Ti2AlN as a function of temperature before and after fitting.
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Figure 12. Schematic illustration for the formation of multi-layered oxide scale during oxidation of TiAlNbCr alloy at 900°C.
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