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Enhanced oxidation resistance of a titanium–based alloy by the addition of boron and the application of electron beam melting
Additive Manufacturing 31 (2020) 100971
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Full Length Article
Enhanced oxidation resistance of a titanium–based alloy by the addition of boron and the application of electron beam melting
T
Yujie Cuia,*, Kenta Aoyagia,*, Yuichiro Koizumia,1, Tadashi Fujiedab, Akihiko Chibaa a b
Institute for Materials Research, Tohoku University, Sendai, 980-8577, Japan Center for Technology Innovation-Materials, Research & Development Group, Hitachi, Ltd., Hitachi, 319-1292, Japan
A R T I C LE I N FO
A B S T R A C T
Keywords: Electron beam powder bed fusion Titanium alloys Oxidation Precipitates
Refined TiB precipitates significantly enhance the oxidation resistance of Ti-6Al-2Sn-4Zr-2Mo-0.1Si-1.0B alloy fabricated by electron beam powder bed fusion (EB-PBF). Refined TiB precipitates in the EB-PBF-built alloy enable finer oxide formation than the larger precipitates in the forged alloy, and the resulting oxidation layers are more compact. Evaporation of scattered B2O3 generated by the refined TiB precipitates in the EB-PBF-built alloy do not significantly accelerate detachment of the oxidation layer from the substrate. However, collective evaporation of B2O3 generated by larger TiB precipitates in the forged alloy accelerate detachment. The oxidation layer on the EB-PBF-fabricated alloy was more stable, preventing further oxidation and improving oxidation resistance.
1. Introduction Electron beam powder bed fusion (EB-PBF) is an additive manufacturing (AM) method that has attracted much attention in recent years due to its significant advantages, which include design freedom and ease of fabricating components with complex shapes that are defect-free and have low residual stresses [1–8]. The EB-PBF process has already been used in the aerospace and medical fields to produce complex parts for which other conventional technologies would be expensive or difficult to apply [9,10]. Titanium (Ti) and its alloys are of particular interest for EB-PBF [11–21]. This is because Ti alloys have broad industrial applications; however, their machining costs are high, and lead times for conventional processing methods are long [22]. Ti-6Al-2Sn-4Zr-2Mo-0.1Si (Ti-6242S) alloy has been used to manufacture jet engine compressors, disks, impellers, and sheet metal components due to its low density, high specific strength, and oxidation resistance [23,24]. However, the mechanical properties and oxidation resistance of Ti-6242S alloy needs to be further improved to meet the design demands of next generation jet engines. The addition of boron can enhance the tensile strength of forged Ti-6242S alloy owing to the formation of TiB precipitates [25]. The strengthening effect of TiB precipitates can be enhanced by EB-PBF manufacturing because refined TiB precipitates are attained with rapid cooling [26]. Although oxidation resistance was evaluated in our previous report [26], differences in
the oxidation behaviour between the forged alloy and the EB-PBF-built Ti-6242S alloy with added boron have not been well characterised. In particular, the contribution of microstructural variations to oxidation resistance needs to be clarified. The oxidation behaviour of Ti-based alloys is closely related to their microstructures. Ti-5.5Al-3.4Sn-3.0Zr-0.7Mo-0.3Si-0.4Nb-0.35Ta alloy fabricated by selective laser melting exhibits better oxidation resistance than the as-cast alloy, which is attributed to its fine martensitic α′ phase, which prevents both the outward diffusion of cations and the influx of oxygen anions during the oxidation process [27]. The surface oxidation of Ti-6Al-4 V alloy has also been shown to be sensitive to microstructural evolution in various EB-PBF-manufactured parts that cool at different rates [28]. In addition, the size and shape of precipitates have been reported to significantly influence oxidation behaviour. Non-steady-state diffusional analysis indicates that fine, homogeneous, and randomly distributed precipitates are more beneficial than coarse precipitates for the exclusive formation of protective oxides in binary multiphase alloys [29]. In short, variations in grain size and precipitate distribution result in tremendously different oxidation behaviour [30,31]. In this study, we aimed to elucidate the differences between the oxidation behaviour of EB-PBF-built and forged Ti-6242S-1.0B alloys caused by different microstructures. We characterised the oxidationlayer constituents and structures in the EB-PBF-built and forged Ti-
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Corresponding authors. E-mail addresses: [email protected] (Y. Cui), [email protected] (K. Aoyagi). 1 Present address: Division of Materials and Manufacturing Science, Graduate School of Engineering, Osaka University, 2-1, Yamadaoka, Suita, Osaka, 565-0871, Japan. https://doi.org/10.1016/j.addma.2019.100971 Received 29 July 2019; Received in revised form 4 November 2019; Accepted 22 November 2019 Available online 24 November 2019 2214-8604/ © 2019 Elsevier B.V. All rights reserved.
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abrasive paper (up to 1000 grit), then polished with a suspension of APA alpha alumina powder (Struers, Ballerup, Denmark). Prior to oxidation, they were ultrasonically cleaned in ethanol and dried with a blower. The samples were placed in ceramic containers to prevent contamination, and isothermal oxidation was performed in air at 800 °C for 2, 4, 8, 16, 24, and 48 h in a standard furnace. After exposure to 800 ± 5 °C for a given duration, the samples were removed and cooled to room temperature in air. The mass gain was measured using a balance (Shimadzu, Japan) with a resolution of 0.01 mg. All fragments collected in the ceramic containers were weighed to ensure that any scales of the oxide layers that may have peeled off during oxidation were included. The oxidation rate was evaluated from the mass gain of the sample using the following equation:
6242S-1.0B alloys by scanning electron microscopy (SEM), X-ray diffractometry (XRD), scanning transmission electron microscopy (STEM), energy dispersive X-ray (EDX) spectroscopy, and high-resolution transmission electron microscopy (HR-TEM). Different factors, especially the role of refined TiB precipitates, for enhancing the oxidation resistance of the EB-PBF-built alloy at high temperatures were investigated. Our findings provide a better understanding of the oxidation behaviour of EB-PBF-fabricated alloys. 2. Experimental 2.1. EB-PBF manufacturing The EB-PBF-built alloy used for this study was manufactured from Ti-6242S-1.0B alloy powder on an Arcam A2X EB-PBF system (Arcam AB, Mölndal, Sweden). The powder was obtained from a high-frequency induction melted electrode (Φ 50 × 500) via a gas atomization process. The average particle size of the powder was 70 μm. The building process was initiated after reaching a preheating temperature of 830 °C under vacuum. The scanning layer thickness was 100 μm, the spot size was 200–300 μm, and the scan speed ranged from 1000 to 2000 mm·s−1. The transverse direction of the beam was rotated by 90° after completing each layer with a line offset of 0.2 mm. The electron beam power was 420 W, and the operating current was 7 mA. Quadrangular prisms with dimensions of 10 × 10 × 40 mm3 were built. The time for the building process was approximately 7 h. The relative density of the EB-PBF-built alloy exceeded 99 %.
Δm = (M1 – M0)/A0,
(1)
where Δm is the mass gain per unit area, M1 is the sample mass after oxidation, M0 is the sample mass before oxidation, and A0 is the surface area of the sample prior to oxidation. Moreover, the parabolic oxidation kinetics can be represented by the relationship [32] Δ m2 = Kpt,
(2)
where t is the exposure time and Kp is the parabolic rate. The macroscopic surface morphologies were initially imaged with a camera (Nikon D5300, Japan). The main constituents of the surface oxidation layer were characterised by XRD analysis. Patterns were collected with an X’Pert Pro MPD diffractometer (Malvern PANalytical, Netherlands) equipped with a Cu Kα (λ = 0.1547 nm) radiation source. The surface morphologies of the oxidation layers were analysed with a S-3400 N scanning electron microscope equipped with a backscattered electron (BSE) detector. To observe the cross-sectional oxidation layers, the samples were mounted in nonconductive epoxy resin at room temperature. After sectioning the samples by EDM, they were ground with SiC abrasive paper (up to 1000 grit), polished with α alumina suspensions, and cleaned in ethanol. Further analysis of the oxidation layer components with HAADF-STEM and HR-TEM was performed on cross-sectional samples prepared by FIB.
2.2. Hot forging To compare the microstructures and oxidation resistances of the EBPBF-built and conventionally wrought Ti-6242S-1.0B alloys, hot forging was performed three times with an arc-plasma melted casting with 30 % reduction at 920 °C. The forged samples were designated ‘as-forged’ alloys. Some of the as-forged samples were firstly treated at 960 °C for 1 h (solution treatment corresponding to the American Society for Testing and Materials standard) and then 830 °C for 7 h followed by air cooling to simulate the thermal history of the post-solidification step in the EB-PBF process. The samples heat treated after forging were designated ‘forged + HT’ alloys.
3. Results 3.1. Initial microstructure
2.3. Microstructure observation
Fig. 1 shows the property diagram of Ti6242S-1.0B alloy calculated by ThemoCalc software. The phases of the Ti6242S-1.0B alloy at room temperature are hexagonally close-packed (hcp) α-Ti, body-centred cubic (bcc) β-Ti, Ti3Al, TiB, and TiZrSi. SEM-BSE images of the samples
The microstructures were analysed prior to oxidation by SEM using a S-3400 N scanning electron microscope (Hitachi, Tokyo, Japan) equipped with an electron probe micro-analyzer (EPMA) detector. Electron backscatter diffraction (EBSD) analysis was carried out for various sample areas (over 300 grains) using the TSL-OIM 5.0 data acquisition software package (TSL Solutions, Japan). The average α phase grain size and the fraction of β-phase can be calculated based on the phase maps obtained by EBSD. TEM samples were prepared by focused ion beam (FIB) milling under a FEI Helios microscope (Thermo Scientific, Waltham, MA). High-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) with dual spherical aberration correction was performed with a Titan3 G2 60–300 TEM at an operating voltage of 300 kV. The thin foils used for HAADF-STEM imaging were cleaned with a 1020 plasma cleaner (Diener Electronic GmbH, Ebhausen, Germany) before they were loaded into the microscope. EDX mapping was carried out during STEM analysis to probe the distribution of elements in the alloys. 2.4. Oxidation tests Electrical discharge machining (EDM) was used to prepare 15 × 15 × 2 mm3 samples of the as-forged, forged + HT, and EB-PBFbuilt alloys for the oxidation tests. All samples were ground with SiC
Fig. 1. Property diagram of Ti6242S-1.0B alloy calculated with ThermoCalc software. 2
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Fig. 2. SEM-BSE images and EPMA elemental maps of (a–e) as-forged, (f–j) forged + heat treated (HT), and (k–o) EB-PBF-built alloys.
Fig. 3. IPF maps, phase maps, and image quality maps of (a–c) as-forged, (d–f) forged + HT, and (g–i) EB-PBF-built alloys.
in Fig. 3. The average grain size was calculated based on the area fraction of different-sized grains. The average α-phase grain sizes of the as-forged, forged + HT, and EB-PBF-built alloys were found to be 2.6 μm ( ± 0.3 μm), 4.2 μm ( ± 0.4 μm), and 3.5 μm ( ± 0.3 μm), respectively. The area fraction of the β-Ti phase was 1.2 % ( ± 0.4 %) in the as-forged alloy, which is similar to that at room temperature calculated from the property diagram under equilibrium solidification (Fig. 1), which was 1.1 %. The area fraction of the β-Ti phase increased
and the corresponding EPMA elemental maps are shown in Fig. 2. The precipitates appear as black spots in the BSE images and were confirmed to be TiB. Large TiB precipitates are randomly distributed in the matrices of the forged alloys, with most of the TiB precipitates in the forged alloys 1–3 μm wide. In contrast, only highly refined TiB precipitates 10–150 nm in size were observed in the EB-PBF-built alloy. The inverse pole figures (IPFs), phase maps, and image quality (IQ) maps of the as-forged, forged + HT, and EB-PBF-built alloys are shown 3
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Fig. 4. (a) Diffraction pattern of TiB precipitate, (b) HAADF-STEM image, and (c–i) STEM-EDX maps of Ti, B, Al, Zr, Sn, Mo, and Si in the EB-PBF-built alloy.
to 6.7 % ( ± 1.2 %) in the forged + HT alloy, which is comparable to that of the EB-PBF-built alloy, which was 8.2 % ( ± 1.4 %). The area fraction of the β-Ti phase was calculated from the property diagram to be around 8.5 % at 830 °C, which is the preheating temperature for the EB-PBF-built alloy and heat treatment temperature for the forged + HT alloy. These results indicate that most of the β-Ti phase at 830 °C remained in the forged + HT and EB-PBF-built alloy owing to the fast cooling rate. The diffraction pattern of the TiB precipitate and STEM-EDX maps of the EB-PBF-built Ti6242S-1.0B alloy are shown in Fig. 4. Based on the diffraction pattern, property diagram, and STEM-EDS results, the nanoscale precipitates with whisker-like morphology were confirmed to be the TiB phase. The EDX maps indicate that the β-Ti phase is rich in molybdenum and boron. The aggregation of boron in the β phase is seemingly supersaturated boron resulting from the high rate of solidification during EB-PBF manufacturing. The cooling rate during EB-PBF manufacturing is estimated to be in the range of 105–106 °C·s−1 [33], which is much higher than conventional casting or forging (in the range of 1–102 °C·s−1 [34,35]). Owing to the faster cooling rate than conventional casting or forging, greater undercooling was generated during EB-PBF manufacturing, which promoted nucleation of TiB precipitates [36]. In addition, the time for the diffusion of boron and the growth of TiB precipitates at elevated temperatures is insufficient owing to short cooling time during EB-PBF manufacturing. Based on the above reasons, the TiB precipitates were significantly refined in the EB-PBF-built alloy compared with the asforged alloy.
Fig. 5. Mass gain per unit area versus the square root of time of the as-forged, forged + HT, and EB-PBF-built [26] alloys at 800 °C.
3.2.2. Microstructural surface morphologies of the oxides The macroscopic surface morphologies of the as-forged and EB-PBFbuilt alloys before and after oxidation at 800 °C for 48 h are shown in Fig. 6a and d. Magnified images of areas A and B in the oxidation layers are shown in Fig. 6b and e, respectively. The oxides in the EB-PBF-built alloy, with an average size of 200 ± 45 nm, are smaller than the oxides in the as-forged alloy (450 ± 60 nm). Cross-sections of the oxidation layers in the as-forged and EB-PBFbuilt alloys exhibit different morphologies, as shown in Fig. 6c and f. The outer oxidation layers detached from the substrate of the as-forged alloy, leaving the inner oxidation layer exposed to the atmosphere. Large TiB precipitates were observed beneath the oxidation layer in the as-forged alloy. In contrast, multiple oxidation layers were observed in the cross-section of the EB-PBF-built alloy that adhered closely to the substrate. These were thinner and more compact than the oxidation layer on the as-forged alloy. The XRD patterns in Fig. 7 verify that the oxides formed on the surfaces of both the as-forged and EB-PBF-built alloys are comprised primarily of rutile TiO2 [37]. The α-Al2O3 [38] and α-Ti [39] phases were also detected in the EB-PBF-built alloy, while the peaks corresponding to α-Al2O3 and α-Ti in the pattern of the as-forged alloy were much less intense. This is because detachment of the oxidation layer reduces the diffusion of Al from the substrate to the surface layer in the as-forged alloy. Consequently, the fraction of Al2O3 in the surface layer of the as-forged alloy is lower than that of the EB-PBF-built alloy. In addition, the extent of oxidation for the EB-PBF-built alloy is less than
3.2. Oxidation behaviour 3.2.1. Oxidation resistance testing The mass gain per unit area (Δm) and the square root of the oxidation time at 800 °C are almost linearly related in the as-forged, forged + HT, and EB-PBF-built alloys, as shown in Fig. 5. The parabolic rate constants of the as-forged, forged + HT, and EB-PBF-built alloys are 0.74, 0.73, and 0.52 mg·cm−2 h-1/2, respectively. The mass gains in the as-forged, forged + HT, and EB-PBF-built alloys after 48 h were found to be 4.43, 4.36, and 2.90 mg·cm−2, respectively. These results indicate that the EB-PBF-built alloy oxidised more slowly and is more resistant to oxidation than the other alloys. The heat treatments with a thermal history simulating that of the post-solidification step in the EB-PBF process had little influence on the oxidation resistance of the as-forged alloy.
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Fig. 6. Surface and cross-sectional morphologies of (a–c) as-forged and (d–f) EB-PBF-built alloys after oxidation testing at 800 °C for 48 h: (a, d) macroscopic surface images, (b, e) magnified SEM images, and (c, f) SEM-BSE images of cross-sectional oxidation layers.
precipitates. This will be further clarified in the discussion. Fig. 10k–o shows that lenticular Al2O3 can nucleate and grow on TiO2 layers in the EB-PBF-built alloy. Consequently, it is possible to form an assemblage of alternating oxidation layers containing TiO2 and Al2O3. It should also be possible for the multiple alternating oxidation layers containing TiO2 and Al2O3 to be formed in the as-forged alloy. However, the detachment of the oxidation layer cut off the diffusion of Al from the substrate to the oxidation layer. Consequently, it is very difficult to grow a thicker Al2O3 layer between the edges of adjacent TiO2 layers in the as-forged alloy. In addition, the lenticular Al2O3 grow along certain directions on the TiO2 layers (Fig. 10m, n). An Al2O3 layer that borders a fine TiO2 layer was observed in the HR-TEM images of the EB-PBF-built alloy, verifying the orientation relationship of the Al2O3 layer and adjacent TiO2 layer (Fig. 11). The orientation of α-Al2O3 relative to rutile TiO2 is well-defined, as shown in Fig. 11c–f. The presence of the [001]TiO2// [10–12]Al2O3 and (010)TiO2//(-2201)Al2O3 orientation relationships is indicative of strong interfacial bonding between rutile TiO2 and αAl2O3; Fig. 12 shows the atomic structures of rutile TiO2 and α-Al2O3. The relationship between their orientations indicates that rutile TiO2 and α-Al2O3 share a well-matched oxygen sublattice, which promotes the growth of α-Al2O3 on rutile TiO2 [40].
Fig. 7. XRD patterns of the oxides formed on the surface of as-forged and EBPBF-built alloys after oxidation at 800 °C for 48 h.
that of the as-forged alloy, resulting in detection of more α-Ti matrix in the EB-PBF-built alloy. 3.2.3. Microstructures of the cross-sectional oxidation layer HAADF images and STEM-EDS maps of the cross-sections of the asforged and EB-PBF-built Ti6242S-1.0B alloys are shown in Fig. 8. The entire oxidation layer of the EB-PBF-built alloy was collected for observation by FIB milling, but only part of the detached oxidation layer was obtained from the as-forged alloy. A transition layer was observed between the oxidation layer and the substrate in the EB-PBF-built alloy. The elemental distributions on the oxidation layers of the as-forged and EB-PBF-built alloys are shown according to their atomic fractions in Fig. 9. The detached, outermost porous oxidation layer of the as-forged alloy was found to be composed primarily of TiO2 as well as some Al2O3. A very thin layer of Al2O3 formed at the boundary with the porous TiO2 layer, as confirmed in the enlarged maps in Fig. 10a–e. Multiple mixed layers beneath the outermost layer are composed mostly of TiO2. The outermost layer of the EB-PBF-built alloy consists mainly of TiO2, as shown in Fig. 10f–j. However, the Al2O3 layer in the interspace between TiO2 layers was thicker than that of the as-forged alloy. A compact assemblage of alternating oxidation layers containing TiO2 and Al2O3 was observed beneath the mixed layer. The formation of a more porous outermost oxidation layer in the as-forged alloy is mainly ascribed to the evaporation of large quantities of B2O3 from TiB
4. Discussion The different oxidation behaviour of the as-forged and EB-PBF-built alloys was elucidated in the present study, and the influence of TiB precipitate size on the oxidation resistance of Ti6242S-1.0B alloys was determined. The formation of a multi-layered structure containing Al2O3 and TiO2 layers was identified. The mechanism for oxidation in the as-forged and EB-PBF-built Ti6242S-1.0B alloys is illustrated schematically in Fig. 13. Outward diffusion of Ti and Al atoms, and diffusion of oxygen into the substrate occur simultaneously [32,41]. Oxygen is adsorbed on the alloy surface at the first stage; oxides then nucleate and grow into compact oxide scales (Fig. 13a and b). The main chemical reactions during the oxidation of Ti-6242S-1.0B alloy take place at high temperatures in an oxidising atmosphere according to Eqs. (3–5):
5
Ti (s) + O (g) → TiO (s)
(3)
TiO (s) + 1/2O2 (g) → TiO2 (s)
(4)
2Al (s) + 3/2O2 (g) → Al2O3 (s)
(5)
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Fig. 8. HAADF images and STEM-EDS maps of the cross-sections of (a–e) as-forged and (f–j) EB-PBF-built alloys.
the surrounding region of Ti, which leads to Al enrichment beneath the TiO2 layer at the second stage. This changes the Al2O3 formation activity to favour Al2O3 growth, resulting in the formation of an Al2O3 layer beneath the TiO2 layer at the third stage, as shown in Fig. 13c. Subsequently, the Al2O3 layers become thicker by depleting the surrounding region of Ti. Due to Ti enrichment beneath the Al2O3 layer, another TiO2 layer forms as more oxygen diffuses into the substrate at the fourth stage, as shown in Fig. 13d. Thus, layers containing TiO2 grow and alternate with Al2O3 layers to form the multi-layered structure of oxide scales. The formation of a multi-layered structure containing Al2O3 and TiO2 layers was also reported in previous research [32,43–47]. For instance, multi-layered structures are observed in Ti6Al-4 V alloy after exposure to the atmosphere at 650−850 °C [32] and in Ti-6Al-2Sn-4Zr-2Mo alloy at 550−650 °C [43]. It was verified that growth of the TiO2 layer leads to depletion of Ti beneath the TiO2 layer, resulting in enhanced formation activity of Al2O3. Similarly, the growth of Al2O3 layer promotes the formation activity of TiO2 beneath the Al2O3 layer. Consequently, multi-layered structure containing Al2O3 and TiO2 layers was formed [32,43]. However, these multi-layered oxide scales can easily detach from the substrate in the as-forged alloy in the present research, as shown in Fig. 13e.
The standard free Gibbs free energies (J/mol) as functions of temperature (T in Kelvin) are represented by [32,42]: ΔGθ (TiO2) = -167600 + 320 T θ
ΔG (Al2O3) = -910000 + 173 T
(6) (7)
θ
ΔG can also be represented by [32]: θ
ΔG (TiO2) = -RTln (aTiO2 /aTi pO2)
(8)
ΔGθ (Al2O3) = -RTln (aAl2O3/(aAl)2(pO2)3/2)
(9)
where a stands for the activity of each phase and pO2 stands for the partial pressure of oxygen. When the unit activities of TiO2, Al2O3, and pO2 are assumed, the minimum values for aTi and aAl to form the respective oxides (TiO2 and Al2O3) can be calculated from Eqs. (6–9). The minimum value of aTi at 800 °C was found to be 2.6 × 10−35, which is much smaller than that of aAl (1.2 × 10-32). Thus, TiO2 is the most likely oxide formed when the external surface of the alloy is exposed to the atmosphere [43,44]. Consequently, TiO2 develops preferentially and forms rapidly on the alloy surface during the second stage of oxidation, as shown in Fig. 13b. With the outward diffusion of Ti and the inward diffusion of oxygen, the TiO2 layers become thicker by depleting
Fig. 9. Atomic fractions of the alloying elements (Ti, Al, O, Zr, Mo, Sn, and Si): (a) from position A to B in Fig. 8 of the as-forged alloy, and (b) from position C to D in Fig. 8 of the EB-PBF-built alloy. 6
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Fig. 10. (a, f, k) Magnified HAADF images and STEM-EDS maps of areas P, Q, and R in Fig. 8 of: (b–e) as-forged and (g–j, l–o) EB-PBF-built alloys.
is ascribed primarily to the influence of the differently sized TiB precipitates. During the oxidation process, TiB precipitates react with O2 according to Eq. 10 [49,50]:
The mismatch between the expansion coefficients of the oxide scale and the substrate results in the formation of cracks in the oxidation layer from the substrate. Moreover, the thermal and growth stresses generated during high-temperature oxidation contribute to scale cracking and spallation [48]. This results in detachment of the oxidation layer, as shown in Fig. 6. However, the extent of oxidation layer detachment from the as-forged and EB-PBF-built alloys differed, which
2TiB (s) + 7/2O2 (g) → 2TiO2 (s) + B2O3
(10)
B2O3 is solid at T < 450 °C and a liquid at T > 450 °C. It has been
Fig. 11. (a, b) HR-TEM images, (c–e) selected areas of electron diffraction patterns, and (f) indexed image of the oxides of the cross-sectional EB-PBF-built alloy. 7
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quantities of scattered B2O3 were generated by the uniformly distributed nanoscale TiB precipitates in the EB-PBF-built alloy. Although evaporation of B2O3 from the EB-PBF-built alloy also reduces the stability of the oxidation layer, the effect was relatively small compared to that observed in the as-forged alloy, as shown in Fig. 13e. The size of the oxides is also thought to significantly affect oxidation resistance. Finer oxides are beneficial for plastic deformation and creep of the oxidation scales [41,51,52]. Cracking and spallation during oxidation were consequently reduced, because scale deformation releases the growth and thermal stresses. TiB precipitates can act as heterogeneous nucleation centres when exposed to oxygen owing to their large surface areas and minimal inter-nuclear distances [53]. The larger number of small TiB precipitates obtained during EB-PBF manufacturing contribute to oxide nucleation, resulting in the formation of finer oxides than those formed in the as-forged alloy (Fig. 6); consequently, finer oxides increase oxidation resistance. In addition, the presence of boron affects the vacancy concentration, which restricts oxygen transport in the oxidation layer [54]. This effect is more pronounced in the EB-PBF-built alloy due to the homogeneous distribution of boron, which results in a slower oxygen diffusion rate [55]. The fraction of Al2O3 in the outermost layer of the EB-PBF-built alloy is higher than in the as-forged alloy. This is possibly because detachment of the oxidation layer in the as-forged alloy halts the diffusion of Al from the substrate to the oxidation layer. The Al2O3 layer is more stable than non-protective TiO2, which also retards the inward diffusion of oxygen and reduces the oxidation rate on the EB-PBF-built alloy [41]. The influences of grain size and the β-phase fraction on oxidation resistance were also considered [45,56,58]. Grain boundaries act as rapid pathways for oxygen, which results in decreased oxidation resistance [56,57]. However, the average α-phase grain sizes in the forged and EB-PBF-built alloys are similar (Fig. 3). Therefore, variation in grain size has little effect on oxidation resistance. Despite heat treatment increasing the percentage of the β-phase from 1.2 % ( ± 0.4 %) to 6.7 % ( ± 1.2 %), the as-forged and forged + HT alloys exhibited similar oxidation resistances (Fig. 5). These results indicate that microstructural evolution during heat treatment has little influence on oxidation resistance. In summary, the enhanced oxidation resistance of the EB-PBF-built
Fig. 12. Atomic structures of rutile TiO2 and α-Al2O3 with the [001]TiO2// [10–12]Al2O3 and (010)TiO2//(-2201)Al2O3 orientation relationships. Red spheres: Ti; green spheres: Al; blue spheres: O (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article).
reported that a considerable amount of B2O3 evaporates at 800 °C [49,50]. TiB precipitates were not detected in the oxidation layer (Figs. 8e and j and 10 e, j, and o), while TiB precipitates can be observed in the substrate below the oxidation layer (Figs. 6c and 8 j). Consequently, we believe that B2O3 in the oxidation layer evaporated. In addition, the supersaturated boron in the β phase may still exist in the oxidation layer of EB-PBF-built alloy after oxidation. The size of TiB precipitates has a significant influence on oxidation behaviour; in the as-forged alloy, large quantities of B2O3 collected around the large TiB precipitates. B2O3 evaporates en masse near 800 °C, which destabilises the oxidation layer and reduces the adhesion force between it and the substrate. The evaporation of B2O3 increases the porosity of the oxidation layer, contributing to the formation of a more porous outermost oxidation layer. Thus, detachment of the oxidation layer from the substrate is accelerated, and new oxidation layers form as the old oxidation layers peel from the surface. This is one reason for the worse oxidation resistance of the as-forged alloy. In contrast, smaller
Fig. 13. Schematic illustration of the oxidation mechanism in as-forged and EB-PBF-built alloys. 8
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alloy is ascribed primarily to the presence of refined TiB precipitates. Evaporation of small quantities of scattered B2O3 do not significantly accelerate detachment of the oxidation layer from the substrate. The finer oxides generated by the refined TiB precipitates decrease the cracking and spallation of oxidation layer. Thus, multiple stable and compact layers are readily formed on the EB-PBF-built alloy, which exhibits better oxidation resistance than the detached oxidation layers of the forged alloy.
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5. Conclusions The TiB precipitates in the EB-PBF-built alloy were significantly refined compared to those in the as-forged alloy, which resulted in the formation of a more stable oxidation layer on the substrate. This is mainly attributed to the effect of TiB precipitate size on the evaporation of B2O3 because the scattered evaporation of small quantities of B2O3 do not significantly accelerate the detachment of the oxidation layer. Owing to the promotion of oxide nucleation by the refined TiB precipitates, finer oxides were formed, which contributed to the formation of multiple compact oxidation layers. Consequently, the resistance of the Ti-6242S-1.0B alloy to oxidation at 800 °C was improved significantly by EB-PBF manufacturing. Acknowledgements The authors would like to thank Yuichiro Hayasaka for help with TEM observations. References [1] D. Herzog, V. Seyda, E. Wycisk, C. Emmelmann, Additive manufacturing of metals, Acta Mater. 117 (2016) 371–392. [2] S. Yoder, S. Morgana, C. Kinzya, E. Barnesa, M. Kirkab, V. Paquitb, P. Nandwanab, A. Plotkowskib, R.R. Dehoffb, S.S. Babua, Characterization of topology optimized Ti-6Al-4V components using electron beam powder bed fusion, Addit. Manuf. 19 (2018) 184–196. [3] T.J. Horn, O.L.A. Harrysson, D.J. Marcellin-Little, H.A. West, B.D.X. Lascelles, R. Aman, Flexural properties of Ti6Al4V rhombic dodecahedron open cellular structures fabricated with electron beam melting, Addit. Manuf. 1 (2014) 2–11. [4] L.E. Murr, Metallurgy of additive manufacturing: examples from electron beam melting, Addit. Manuf. 5 (2015) 40–53. [5] P. Wang, X. Tan, C. He, M. Nai, R. Huang, S. Tor, J. Wei, Scanning optical microscopy for porosity quantification of additively manufactured components, Addit. Manuf. 21 (2018) 350–358. [6] N. Béraud, F. Vignat, F. Villeneuve, R. Dendievel, Improving dimensional accuracy in EBM using beam characterization and trajectory optimization, Addit. Manuf. 14 (2017) 1–6. [7] J.K. Algardh, T. Hornb, H. Westb, R. Amanb, A. Snisc, H. Engqvistd, J. Lausmaaa, O. Harrysson, Thickness dependency of mechanical properties for thin-walled titanium parts manufactured by Electron beam Melting (EBM), Addit. Manuf. 12 (2016) 45–50. [8] C.A. Terrazas, L.E. Murr, D. Bermudez, E. Arrieta, D.A. Roberson, R.B. Wicker, Microstructure and mechanical properties of Ti-6Al-4V-5% hydroxyapatite composite fabricated using electron beam powder bed fusion, J. Mater. Sci. Technol. 35 (2019) 309–321. [9] M. Galati, L. Iuliano, A literature review of powder-based electron beam melting focusing on numerical simulations, Addit. Manuf. 19 (2018) 1–20. [10] I. Eldesouky, O. Harrysson, H. West, H. Elhofy, Electron beam melted scaffolds for orthopedic applications, Addit. Manuf. 17 (2017) 169–175. [11] M. Todai, T. Nakanoa, T. Liua, H.Y. Yasudaa, K. Hagiharab, K. Choa, M. Uedac, M. Takeyamad, Effect of building direction on the microstructure and tensile properties of Ti-48Al-2Cr-2Nb alloy additively manufactured by electron beam melting, Addit. Manuf. 13 (2017) 61–70. [12] W.A. Grell, E. Solis-Ramosa, E. Clarka, E. Luconc, E.J. Garboczic, P.K. Predeckia, Z. Loftusb, M. Kumosaa, Effect of powder oxidation on the impact toughness of electron beam melting Ti-6Al-4V, Addit. Manuf. 17 (2017) 123–134. [13] V.V. Popov, A. Katz-Demyanetz, A. Garkun, M. Bamberger, The effect of powder recycling on the mechanical properties and microstructure of electron beam melted Ti-6Al-4 V specimens, Addit. Manuf. 22 (2018) 834–843. [14] H. Galarraga, D.A. Lados, R.R. Dehoff, M.M. Kirka, P. Nandwana, Effects of the microstructure and porosity on properties of Ti-6Al-4V ELI alloy fabricated by electron beam melting (EBM), Addit. Manuf. 10 (2016) 47–57. [15] W. Everhart, J. Dinardo, Process driven strengthening mechanisms in electron beam melted Ti-6Al-4V, Addit. Manuf. 24 (2018) 440–445. [16] P. Wang, X. Tan, M.L.S. Nai, S.B. Tor, J. Wei, Spatial and geometrical-based characterization of microstructure and microhardness for an electron beam melted Ti-
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material, Oxid. Met. 57 (2002) 99–114. [52] V. Parry, M.T. Tran, Y. Wouters, Detection of breakaway oxidation with acoustic emission during zirconium oxide scale growth, Oxid. Met. 79 (2013) 279–288. [53] Y. Qin, D. Zhang, W. Lu, J. Qin, D. Xu, Oxidation behavior of in situ synthesized TiB/Ti-Al composite, J. Mater. Sci. 24 (2005) 6553–6558. [54] D.A. Brice, P. Samimi, I. Ghamarian, Y. Liu, R.M. Brice, R.F. Reidy, J.D. Cotton, M.J. Kaufman, P.C. Collins, Oxidation behavior and microstructural decomposition of Ti-6Al-4V and Ti-6Al-4V-1B sheet, Corros. Sci. 112 (2016) 338–346. [55] L. Su, L. Jia, J. Weng, Z. Hong, C. Zhou, H. Zhang, Improvement in the oxidation
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Full Length Article
Enhanced oxidation resistance of a titanium–based alloy by the addition of boron and the application of electron beam melting
T
Yujie Cuia,*, Kenta Aoyagia,*, Yuichiro Koizumia,1, Tadashi Fujiedab, Akihiko Chibaa a b
Institute for Materials Research, Tohoku University, Sendai, 980-8577, Japan Center for Technology Innovation-Materials, Research & Development Group, Hitachi, Ltd., Hitachi, 319-1292, Japan
A R T I C LE I N FO
A B S T R A C T
Keywords: Electron beam powder bed fusion Titanium alloys Oxidation Precipitates
Refined TiB precipitates significantly enhance the oxidation resistance of Ti-6Al-2Sn-4Zr-2Mo-0.1Si-1.0B alloy fabricated by electron beam powder bed fusion (EB-PBF). Refined TiB precipitates in the EB-PBF-built alloy enable finer oxide formation than the larger precipitates in the forged alloy, and the resulting oxidation layers are more compact. Evaporation of scattered B2O3 generated by the refined TiB precipitates in the EB-PBF-built alloy do not significantly accelerate detachment of the oxidation layer from the substrate. However, collective evaporation of B2O3 generated by larger TiB precipitates in the forged alloy accelerate detachment. The oxidation layer on the EB-PBF-fabricated alloy was more stable, preventing further oxidation and improving oxidation resistance.
1. Introduction Electron beam powder bed fusion (EB-PBF) is an additive manufacturing (AM) method that has attracted much attention in recent years due to its significant advantages, which include design freedom and ease of fabricating components with complex shapes that are defect-free and have low residual stresses [1–8]. The EB-PBF process has already been used in the aerospace and medical fields to produce complex parts for which other conventional technologies would be expensive or difficult to apply [9,10]. Titanium (Ti) and its alloys are of particular interest for EB-PBF [11–21]. This is because Ti alloys have broad industrial applications; however, their machining costs are high, and lead times for conventional processing methods are long [22]. Ti-6Al-2Sn-4Zr-2Mo-0.1Si (Ti-6242S) alloy has been used to manufacture jet engine compressors, disks, impellers, and sheet metal components due to its low density, high specific strength, and oxidation resistance [23,24]. However, the mechanical properties and oxidation resistance of Ti-6242S alloy needs to be further improved to meet the design demands of next generation jet engines. The addition of boron can enhance the tensile strength of forged Ti-6242S alloy owing to the formation of TiB precipitates [25]. The strengthening effect of TiB precipitates can be enhanced by EB-PBF manufacturing because refined TiB precipitates are attained with rapid cooling [26]. Although oxidation resistance was evaluated in our previous report [26], differences in
the oxidation behaviour between the forged alloy and the EB-PBF-built Ti-6242S alloy with added boron have not been well characterised. In particular, the contribution of microstructural variations to oxidation resistance needs to be clarified. The oxidation behaviour of Ti-based alloys is closely related to their microstructures. Ti-5.5Al-3.4Sn-3.0Zr-0.7Mo-0.3Si-0.4Nb-0.35Ta alloy fabricated by selective laser melting exhibits better oxidation resistance than the as-cast alloy, which is attributed to its fine martensitic α′ phase, which prevents both the outward diffusion of cations and the influx of oxygen anions during the oxidation process [27]. The surface oxidation of Ti-6Al-4 V alloy has also been shown to be sensitive to microstructural evolution in various EB-PBF-manufactured parts that cool at different rates [28]. In addition, the size and shape of precipitates have been reported to significantly influence oxidation behaviour. Non-steady-state diffusional analysis indicates that fine, homogeneous, and randomly distributed precipitates are more beneficial than coarse precipitates for the exclusive formation of protective oxides in binary multiphase alloys [29]. In short, variations in grain size and precipitate distribution result in tremendously different oxidation behaviour [30,31]. In this study, we aimed to elucidate the differences between the oxidation behaviour of EB-PBF-built and forged Ti-6242S-1.0B alloys caused by different microstructures. We characterised the oxidationlayer constituents and structures in the EB-PBF-built and forged Ti-
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Corresponding authors. E-mail addresses: [email protected] (Y. Cui), [email protected] (K. Aoyagi). 1 Present address: Division of Materials and Manufacturing Science, Graduate School of Engineering, Osaka University, 2-1, Yamadaoka, Suita, Osaka, 565-0871, Japan. https://doi.org/10.1016/j.addma.2019.100971 Received 29 July 2019; Received in revised form 4 November 2019; Accepted 22 November 2019 Available online 24 November 2019 2214-8604/ © 2019 Elsevier B.V. All rights reserved.
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abrasive paper (up to 1000 grit), then polished with a suspension of APA alpha alumina powder (Struers, Ballerup, Denmark). Prior to oxidation, they were ultrasonically cleaned in ethanol and dried with a blower. The samples were placed in ceramic containers to prevent contamination, and isothermal oxidation was performed in air at 800 °C for 2, 4, 8, 16, 24, and 48 h in a standard furnace. After exposure to 800 ± 5 °C for a given duration, the samples were removed and cooled to room temperature in air. The mass gain was measured using a balance (Shimadzu, Japan) with a resolution of 0.01 mg. All fragments collected in the ceramic containers were weighed to ensure that any scales of the oxide layers that may have peeled off during oxidation were included. The oxidation rate was evaluated from the mass gain of the sample using the following equation:
6242S-1.0B alloys by scanning electron microscopy (SEM), X-ray diffractometry (XRD), scanning transmission electron microscopy (STEM), energy dispersive X-ray (EDX) spectroscopy, and high-resolution transmission electron microscopy (HR-TEM). Different factors, especially the role of refined TiB precipitates, for enhancing the oxidation resistance of the EB-PBF-built alloy at high temperatures were investigated. Our findings provide a better understanding of the oxidation behaviour of EB-PBF-fabricated alloys. 2. Experimental 2.1. EB-PBF manufacturing The EB-PBF-built alloy used for this study was manufactured from Ti-6242S-1.0B alloy powder on an Arcam A2X EB-PBF system (Arcam AB, Mölndal, Sweden). The powder was obtained from a high-frequency induction melted electrode (Φ 50 × 500) via a gas atomization process. The average particle size of the powder was 70 μm. The building process was initiated after reaching a preheating temperature of 830 °C under vacuum. The scanning layer thickness was 100 μm, the spot size was 200–300 μm, and the scan speed ranged from 1000 to 2000 mm·s−1. The transverse direction of the beam was rotated by 90° after completing each layer with a line offset of 0.2 mm. The electron beam power was 420 W, and the operating current was 7 mA. Quadrangular prisms with dimensions of 10 × 10 × 40 mm3 were built. The time for the building process was approximately 7 h. The relative density of the EB-PBF-built alloy exceeded 99 %.
Δm = (M1 – M0)/A0,
(1)
where Δm is the mass gain per unit area, M1 is the sample mass after oxidation, M0 is the sample mass before oxidation, and A0 is the surface area of the sample prior to oxidation. Moreover, the parabolic oxidation kinetics can be represented by the relationship [32] Δ m2 = Kpt,
(2)
where t is the exposure time and Kp is the parabolic rate. The macroscopic surface morphologies were initially imaged with a camera (Nikon D5300, Japan). The main constituents of the surface oxidation layer were characterised by XRD analysis. Patterns were collected with an X’Pert Pro MPD diffractometer (Malvern PANalytical, Netherlands) equipped with a Cu Kα (λ = 0.1547 nm) radiation source. The surface morphologies of the oxidation layers were analysed with a S-3400 N scanning electron microscope equipped with a backscattered electron (BSE) detector. To observe the cross-sectional oxidation layers, the samples were mounted in nonconductive epoxy resin at room temperature. After sectioning the samples by EDM, they were ground with SiC abrasive paper (up to 1000 grit), polished with α alumina suspensions, and cleaned in ethanol. Further analysis of the oxidation layer components with HAADF-STEM and HR-TEM was performed on cross-sectional samples prepared by FIB.
2.2. Hot forging To compare the microstructures and oxidation resistances of the EBPBF-built and conventionally wrought Ti-6242S-1.0B alloys, hot forging was performed three times with an arc-plasma melted casting with 30 % reduction at 920 °C. The forged samples were designated ‘as-forged’ alloys. Some of the as-forged samples were firstly treated at 960 °C for 1 h (solution treatment corresponding to the American Society for Testing and Materials standard) and then 830 °C for 7 h followed by air cooling to simulate the thermal history of the post-solidification step in the EB-PBF process. The samples heat treated after forging were designated ‘forged + HT’ alloys.
3. Results 3.1. Initial microstructure
2.3. Microstructure observation
Fig. 1 shows the property diagram of Ti6242S-1.0B alloy calculated by ThemoCalc software. The phases of the Ti6242S-1.0B alloy at room temperature are hexagonally close-packed (hcp) α-Ti, body-centred cubic (bcc) β-Ti, Ti3Al, TiB, and TiZrSi. SEM-BSE images of the samples
The microstructures were analysed prior to oxidation by SEM using a S-3400 N scanning electron microscope (Hitachi, Tokyo, Japan) equipped with an electron probe micro-analyzer (EPMA) detector. Electron backscatter diffraction (EBSD) analysis was carried out for various sample areas (over 300 grains) using the TSL-OIM 5.0 data acquisition software package (TSL Solutions, Japan). The average α phase grain size and the fraction of β-phase can be calculated based on the phase maps obtained by EBSD. TEM samples were prepared by focused ion beam (FIB) milling under a FEI Helios microscope (Thermo Scientific, Waltham, MA). High-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) with dual spherical aberration correction was performed with a Titan3 G2 60–300 TEM at an operating voltage of 300 kV. The thin foils used for HAADF-STEM imaging were cleaned with a 1020 plasma cleaner (Diener Electronic GmbH, Ebhausen, Germany) before they were loaded into the microscope. EDX mapping was carried out during STEM analysis to probe the distribution of elements in the alloys. 2.4. Oxidation tests Electrical discharge machining (EDM) was used to prepare 15 × 15 × 2 mm3 samples of the as-forged, forged + HT, and EB-PBFbuilt alloys for the oxidation tests. All samples were ground with SiC
Fig. 1. Property diagram of Ti6242S-1.0B alloy calculated with ThermoCalc software. 2
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Fig. 2. SEM-BSE images and EPMA elemental maps of (a–e) as-forged, (f–j) forged + heat treated (HT), and (k–o) EB-PBF-built alloys.
Fig. 3. IPF maps, phase maps, and image quality maps of (a–c) as-forged, (d–f) forged + HT, and (g–i) EB-PBF-built alloys.
in Fig. 3. The average grain size was calculated based on the area fraction of different-sized grains. The average α-phase grain sizes of the as-forged, forged + HT, and EB-PBF-built alloys were found to be 2.6 μm ( ± 0.3 μm), 4.2 μm ( ± 0.4 μm), and 3.5 μm ( ± 0.3 μm), respectively. The area fraction of the β-Ti phase was 1.2 % ( ± 0.4 %) in the as-forged alloy, which is similar to that at room temperature calculated from the property diagram under equilibrium solidification (Fig. 1), which was 1.1 %. The area fraction of the β-Ti phase increased
and the corresponding EPMA elemental maps are shown in Fig. 2. The precipitates appear as black spots in the BSE images and were confirmed to be TiB. Large TiB precipitates are randomly distributed in the matrices of the forged alloys, with most of the TiB precipitates in the forged alloys 1–3 μm wide. In contrast, only highly refined TiB precipitates 10–150 nm in size were observed in the EB-PBF-built alloy. The inverse pole figures (IPFs), phase maps, and image quality (IQ) maps of the as-forged, forged + HT, and EB-PBF-built alloys are shown 3
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Fig. 4. (a) Diffraction pattern of TiB precipitate, (b) HAADF-STEM image, and (c–i) STEM-EDX maps of Ti, B, Al, Zr, Sn, Mo, and Si in the EB-PBF-built alloy.
to 6.7 % ( ± 1.2 %) in the forged + HT alloy, which is comparable to that of the EB-PBF-built alloy, which was 8.2 % ( ± 1.4 %). The area fraction of the β-Ti phase was calculated from the property diagram to be around 8.5 % at 830 °C, which is the preheating temperature for the EB-PBF-built alloy and heat treatment temperature for the forged + HT alloy. These results indicate that most of the β-Ti phase at 830 °C remained in the forged + HT and EB-PBF-built alloy owing to the fast cooling rate. The diffraction pattern of the TiB precipitate and STEM-EDX maps of the EB-PBF-built Ti6242S-1.0B alloy are shown in Fig. 4. Based on the diffraction pattern, property diagram, and STEM-EDS results, the nanoscale precipitates with whisker-like morphology were confirmed to be the TiB phase. The EDX maps indicate that the β-Ti phase is rich in molybdenum and boron. The aggregation of boron in the β phase is seemingly supersaturated boron resulting from the high rate of solidification during EB-PBF manufacturing. The cooling rate during EB-PBF manufacturing is estimated to be in the range of 105–106 °C·s−1 [33], which is much higher than conventional casting or forging (in the range of 1–102 °C·s−1 [34,35]). Owing to the faster cooling rate than conventional casting or forging, greater undercooling was generated during EB-PBF manufacturing, which promoted nucleation of TiB precipitates [36]. In addition, the time for the diffusion of boron and the growth of TiB precipitates at elevated temperatures is insufficient owing to short cooling time during EB-PBF manufacturing. Based on the above reasons, the TiB precipitates were significantly refined in the EB-PBF-built alloy compared with the asforged alloy.
Fig. 5. Mass gain per unit area versus the square root of time of the as-forged, forged + HT, and EB-PBF-built [26] alloys at 800 °C.
3.2.2. Microstructural surface morphologies of the oxides The macroscopic surface morphologies of the as-forged and EB-PBFbuilt alloys before and after oxidation at 800 °C for 48 h are shown in Fig. 6a and d. Magnified images of areas A and B in the oxidation layers are shown in Fig. 6b and e, respectively. The oxides in the EB-PBF-built alloy, with an average size of 200 ± 45 nm, are smaller than the oxides in the as-forged alloy (450 ± 60 nm). Cross-sections of the oxidation layers in the as-forged and EB-PBFbuilt alloys exhibit different morphologies, as shown in Fig. 6c and f. The outer oxidation layers detached from the substrate of the as-forged alloy, leaving the inner oxidation layer exposed to the atmosphere. Large TiB precipitates were observed beneath the oxidation layer in the as-forged alloy. In contrast, multiple oxidation layers were observed in the cross-section of the EB-PBF-built alloy that adhered closely to the substrate. These were thinner and more compact than the oxidation layer on the as-forged alloy. The XRD patterns in Fig. 7 verify that the oxides formed on the surfaces of both the as-forged and EB-PBF-built alloys are comprised primarily of rutile TiO2 [37]. The α-Al2O3 [38] and α-Ti [39] phases were also detected in the EB-PBF-built alloy, while the peaks corresponding to α-Al2O3 and α-Ti in the pattern of the as-forged alloy were much less intense. This is because detachment of the oxidation layer reduces the diffusion of Al from the substrate to the surface layer in the as-forged alloy. Consequently, the fraction of Al2O3 in the surface layer of the as-forged alloy is lower than that of the EB-PBF-built alloy. In addition, the extent of oxidation for the EB-PBF-built alloy is less than
3.2. Oxidation behaviour 3.2.1. Oxidation resistance testing The mass gain per unit area (Δm) and the square root of the oxidation time at 800 °C are almost linearly related in the as-forged, forged + HT, and EB-PBF-built alloys, as shown in Fig. 5. The parabolic rate constants of the as-forged, forged + HT, and EB-PBF-built alloys are 0.74, 0.73, and 0.52 mg·cm−2 h-1/2, respectively. The mass gains in the as-forged, forged + HT, and EB-PBF-built alloys after 48 h were found to be 4.43, 4.36, and 2.90 mg·cm−2, respectively. These results indicate that the EB-PBF-built alloy oxidised more slowly and is more resistant to oxidation than the other alloys. The heat treatments with a thermal history simulating that of the post-solidification step in the EB-PBF process had little influence on the oxidation resistance of the as-forged alloy.
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Fig. 6. Surface and cross-sectional morphologies of (a–c) as-forged and (d–f) EB-PBF-built alloys after oxidation testing at 800 °C for 48 h: (a, d) macroscopic surface images, (b, e) magnified SEM images, and (c, f) SEM-BSE images of cross-sectional oxidation layers.
precipitates. This will be further clarified in the discussion. Fig. 10k–o shows that lenticular Al2O3 can nucleate and grow on TiO2 layers in the EB-PBF-built alloy. Consequently, it is possible to form an assemblage of alternating oxidation layers containing TiO2 and Al2O3. It should also be possible for the multiple alternating oxidation layers containing TiO2 and Al2O3 to be formed in the as-forged alloy. However, the detachment of the oxidation layer cut off the diffusion of Al from the substrate to the oxidation layer. Consequently, it is very difficult to grow a thicker Al2O3 layer between the edges of adjacent TiO2 layers in the as-forged alloy. In addition, the lenticular Al2O3 grow along certain directions on the TiO2 layers (Fig. 10m, n). An Al2O3 layer that borders a fine TiO2 layer was observed in the HR-TEM images of the EB-PBF-built alloy, verifying the orientation relationship of the Al2O3 layer and adjacent TiO2 layer (Fig. 11). The orientation of α-Al2O3 relative to rutile TiO2 is well-defined, as shown in Fig. 11c–f. The presence of the [001]TiO2// [10–12]Al2O3 and (010)TiO2//(-2201)Al2O3 orientation relationships is indicative of strong interfacial bonding between rutile TiO2 and αAl2O3; Fig. 12 shows the atomic structures of rutile TiO2 and α-Al2O3. The relationship between their orientations indicates that rutile TiO2 and α-Al2O3 share a well-matched oxygen sublattice, which promotes the growth of α-Al2O3 on rutile TiO2 [40].
Fig. 7. XRD patterns of the oxides formed on the surface of as-forged and EBPBF-built alloys after oxidation at 800 °C for 48 h.
that of the as-forged alloy, resulting in detection of more α-Ti matrix in the EB-PBF-built alloy. 3.2.3. Microstructures of the cross-sectional oxidation layer HAADF images and STEM-EDS maps of the cross-sections of the asforged and EB-PBF-built Ti6242S-1.0B alloys are shown in Fig. 8. The entire oxidation layer of the EB-PBF-built alloy was collected for observation by FIB milling, but only part of the detached oxidation layer was obtained from the as-forged alloy. A transition layer was observed between the oxidation layer and the substrate in the EB-PBF-built alloy. The elemental distributions on the oxidation layers of the as-forged and EB-PBF-built alloys are shown according to their atomic fractions in Fig. 9. The detached, outermost porous oxidation layer of the as-forged alloy was found to be composed primarily of TiO2 as well as some Al2O3. A very thin layer of Al2O3 formed at the boundary with the porous TiO2 layer, as confirmed in the enlarged maps in Fig. 10a–e. Multiple mixed layers beneath the outermost layer are composed mostly of TiO2. The outermost layer of the EB-PBF-built alloy consists mainly of TiO2, as shown in Fig. 10f–j. However, the Al2O3 layer in the interspace between TiO2 layers was thicker than that of the as-forged alloy. A compact assemblage of alternating oxidation layers containing TiO2 and Al2O3 was observed beneath the mixed layer. The formation of a more porous outermost oxidation layer in the as-forged alloy is mainly ascribed to the evaporation of large quantities of B2O3 from TiB
4. Discussion The different oxidation behaviour of the as-forged and EB-PBF-built alloys was elucidated in the present study, and the influence of TiB precipitate size on the oxidation resistance of Ti6242S-1.0B alloys was determined. The formation of a multi-layered structure containing Al2O3 and TiO2 layers was identified. The mechanism for oxidation in the as-forged and EB-PBF-built Ti6242S-1.0B alloys is illustrated schematically in Fig. 13. Outward diffusion of Ti and Al atoms, and diffusion of oxygen into the substrate occur simultaneously [32,41]. Oxygen is adsorbed on the alloy surface at the first stage; oxides then nucleate and grow into compact oxide scales (Fig. 13a and b). The main chemical reactions during the oxidation of Ti-6242S-1.0B alloy take place at high temperatures in an oxidising atmosphere according to Eqs. (3–5):
5
Ti (s) + O (g) → TiO (s)
(3)
TiO (s) + 1/2O2 (g) → TiO2 (s)
(4)
2Al (s) + 3/2O2 (g) → Al2O3 (s)
(5)
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Fig. 8. HAADF images and STEM-EDS maps of the cross-sections of (a–e) as-forged and (f–j) EB-PBF-built alloys.
the surrounding region of Ti, which leads to Al enrichment beneath the TiO2 layer at the second stage. This changes the Al2O3 formation activity to favour Al2O3 growth, resulting in the formation of an Al2O3 layer beneath the TiO2 layer at the third stage, as shown in Fig. 13c. Subsequently, the Al2O3 layers become thicker by depleting the surrounding region of Ti. Due to Ti enrichment beneath the Al2O3 layer, another TiO2 layer forms as more oxygen diffuses into the substrate at the fourth stage, as shown in Fig. 13d. Thus, layers containing TiO2 grow and alternate with Al2O3 layers to form the multi-layered structure of oxide scales. The formation of a multi-layered structure containing Al2O3 and TiO2 layers was also reported in previous research [32,43–47]. For instance, multi-layered structures are observed in Ti6Al-4 V alloy after exposure to the atmosphere at 650−850 °C [32] and in Ti-6Al-2Sn-4Zr-2Mo alloy at 550−650 °C [43]. It was verified that growth of the TiO2 layer leads to depletion of Ti beneath the TiO2 layer, resulting in enhanced formation activity of Al2O3. Similarly, the growth of Al2O3 layer promotes the formation activity of TiO2 beneath the Al2O3 layer. Consequently, multi-layered structure containing Al2O3 and TiO2 layers was formed [32,43]. However, these multi-layered oxide scales can easily detach from the substrate in the as-forged alloy in the present research, as shown in Fig. 13e.
The standard free Gibbs free energies (J/mol) as functions of temperature (T in Kelvin) are represented by [32,42]: ΔGθ (TiO2) = -167600 + 320 T θ
ΔG (Al2O3) = -910000 + 173 T
(6) (7)
θ
ΔG can also be represented by [32]: θ
ΔG (TiO2) = -RTln (aTiO2 /aTi pO2)
(8)
ΔGθ (Al2O3) = -RTln (aAl2O3/(aAl)2(pO2)3/2)
(9)
where a stands for the activity of each phase and pO2 stands for the partial pressure of oxygen. When the unit activities of TiO2, Al2O3, and pO2 are assumed, the minimum values for aTi and aAl to form the respective oxides (TiO2 and Al2O3) can be calculated from Eqs. (6–9). The minimum value of aTi at 800 °C was found to be 2.6 × 10−35, which is much smaller than that of aAl (1.2 × 10-32). Thus, TiO2 is the most likely oxide formed when the external surface of the alloy is exposed to the atmosphere [43,44]. Consequently, TiO2 develops preferentially and forms rapidly on the alloy surface during the second stage of oxidation, as shown in Fig. 13b. With the outward diffusion of Ti and the inward diffusion of oxygen, the TiO2 layers become thicker by depleting
Fig. 9. Atomic fractions of the alloying elements (Ti, Al, O, Zr, Mo, Sn, and Si): (a) from position A to B in Fig. 8 of the as-forged alloy, and (b) from position C to D in Fig. 8 of the EB-PBF-built alloy. 6
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Fig. 10. (a, f, k) Magnified HAADF images and STEM-EDS maps of areas P, Q, and R in Fig. 8 of: (b–e) as-forged and (g–j, l–o) EB-PBF-built alloys.
is ascribed primarily to the influence of the differently sized TiB precipitates. During the oxidation process, TiB precipitates react with O2 according to Eq. 10 [49,50]:
The mismatch between the expansion coefficients of the oxide scale and the substrate results in the formation of cracks in the oxidation layer from the substrate. Moreover, the thermal and growth stresses generated during high-temperature oxidation contribute to scale cracking and spallation [48]. This results in detachment of the oxidation layer, as shown in Fig. 6. However, the extent of oxidation layer detachment from the as-forged and EB-PBF-built alloys differed, which
2TiB (s) + 7/2O2 (g) → 2TiO2 (s) + B2O3
(10)
B2O3 is solid at T < 450 °C and a liquid at T > 450 °C. It has been
Fig. 11. (a, b) HR-TEM images, (c–e) selected areas of electron diffraction patterns, and (f) indexed image of the oxides of the cross-sectional EB-PBF-built alloy. 7
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quantities of scattered B2O3 were generated by the uniformly distributed nanoscale TiB precipitates in the EB-PBF-built alloy. Although evaporation of B2O3 from the EB-PBF-built alloy also reduces the stability of the oxidation layer, the effect was relatively small compared to that observed in the as-forged alloy, as shown in Fig. 13e. The size of the oxides is also thought to significantly affect oxidation resistance. Finer oxides are beneficial for plastic deformation and creep of the oxidation scales [41,51,52]. Cracking and spallation during oxidation were consequently reduced, because scale deformation releases the growth and thermal stresses. TiB precipitates can act as heterogeneous nucleation centres when exposed to oxygen owing to their large surface areas and minimal inter-nuclear distances [53]. The larger number of small TiB precipitates obtained during EB-PBF manufacturing contribute to oxide nucleation, resulting in the formation of finer oxides than those formed in the as-forged alloy (Fig. 6); consequently, finer oxides increase oxidation resistance. In addition, the presence of boron affects the vacancy concentration, which restricts oxygen transport in the oxidation layer [54]. This effect is more pronounced in the EB-PBF-built alloy due to the homogeneous distribution of boron, which results in a slower oxygen diffusion rate [55]. The fraction of Al2O3 in the outermost layer of the EB-PBF-built alloy is higher than in the as-forged alloy. This is possibly because detachment of the oxidation layer in the as-forged alloy halts the diffusion of Al from the substrate to the oxidation layer. The Al2O3 layer is more stable than non-protective TiO2, which also retards the inward diffusion of oxygen and reduces the oxidation rate on the EB-PBF-built alloy [41]. The influences of grain size and the β-phase fraction on oxidation resistance were also considered [45,56,58]. Grain boundaries act as rapid pathways for oxygen, which results in decreased oxidation resistance [56,57]. However, the average α-phase grain sizes in the forged and EB-PBF-built alloys are similar (Fig. 3). Therefore, variation in grain size has little effect on oxidation resistance. Despite heat treatment increasing the percentage of the β-phase from 1.2 % ( ± 0.4 %) to 6.7 % ( ± 1.2 %), the as-forged and forged + HT alloys exhibited similar oxidation resistances (Fig. 5). These results indicate that microstructural evolution during heat treatment has little influence on oxidation resistance. In summary, the enhanced oxidation resistance of the EB-PBF-built
Fig. 12. Atomic structures of rutile TiO2 and α-Al2O3 with the [001]TiO2// [10–12]Al2O3 and (010)TiO2//(-2201)Al2O3 orientation relationships. Red spheres: Ti; green spheres: Al; blue spheres: O (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article).
reported that a considerable amount of B2O3 evaporates at 800 °C [49,50]. TiB precipitates were not detected in the oxidation layer (Figs. 8e and j and 10 e, j, and o), while TiB precipitates can be observed in the substrate below the oxidation layer (Figs. 6c and 8 j). Consequently, we believe that B2O3 in the oxidation layer evaporated. In addition, the supersaturated boron in the β phase may still exist in the oxidation layer of EB-PBF-built alloy after oxidation. The size of TiB precipitates has a significant influence on oxidation behaviour; in the as-forged alloy, large quantities of B2O3 collected around the large TiB precipitates. B2O3 evaporates en masse near 800 °C, which destabilises the oxidation layer and reduces the adhesion force between it and the substrate. The evaporation of B2O3 increases the porosity of the oxidation layer, contributing to the formation of a more porous outermost oxidation layer. Thus, detachment of the oxidation layer from the substrate is accelerated, and new oxidation layers form as the old oxidation layers peel from the surface. This is one reason for the worse oxidation resistance of the as-forged alloy. In contrast, smaller
Fig. 13. Schematic illustration of the oxidation mechanism in as-forged and EB-PBF-built alloys. 8
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alloy is ascribed primarily to the presence of refined TiB precipitates. Evaporation of small quantities of scattered B2O3 do not significantly accelerate detachment of the oxidation layer from the substrate. The finer oxides generated by the refined TiB precipitates decrease the cracking and spallation of oxidation layer. Thus, multiple stable and compact layers are readily formed on the EB-PBF-built alloy, which exhibits better oxidation resistance than the detached oxidation layers of the forged alloy.
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