| Cite this article as: |
Yang Liu, Yufeng Liu, Sha Zhang, Lin Zhang, Peng Zhang, Shaorong Zhang, Na Liu, Zhou Li, and Xuanhui Qu, Structure characterization of the oxide film on FGH96 superalloy powders with various oxidation degrees, Int. J. Miner. Metall. Mater., 31(2024), No. 9, pp.2037-2047. https://dx.doi.org/10.1007/s12613-024-2823-x
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The aerospace industry demands high-performance turbine blades made from Ni-based superalloys that meet strict requirements for uniform microstructure and thermal working performance [1]. Traditional casting and forging methods are insufficient in addressing these limitations [2–4]. Herein, powder metallurgy superalloys have been developed as an alternative due to their superior mechanical properties at high temperatures and uniform microstructure [5–9]. However, the formation of carbides and oxides at the prior particle boundaries (PPBs) during the solidification of the alloy powders can negatively impact the mechanical properties of the superalloys [10–11].
Numerous works have focused on the evolution of PPBs in Ni-based superalloys, especially the carbides at the PPBs [12–15]. Research suggested that the formation of carbides on the powder surface is related to the segregation of the C and alloy elements. The semi-stable carbides would decompose during consolidation, while stable carbides are inherited in the subsequent process [10]. However, in other researches [16–18], the powder surface of FGH96 was suggested as a composed of various oxide particles without carbides. Since the different surface structure of the alloy powders could lead to the different oxidation behaviors, it is necessary to study the surface structure of the FGH96 alloy powders.
Besides, the negative effect of oxides on the mechanical properties of superalloys cannot be ignored. For example, researches [19−20] showed that the oxide particles at PPBs would accelerate the intergranular fatigue crack growth due to their brittle nature and low fracture toughness. Besides, the oxygen content could also affect the formation of twin boundaries, which would influence the high temperature properties of the Ni-based superalloys [21]. Nevertheless, understanding the characterization of oxygen existing on the surface of the powder surface and its subsequent conversion into oxides on the grain boundaries of the PPBs remains ambiguous.
The research on oxide film is also limited by the difficulties in characterization caused by low oxygen content on the powder surface. To address this issue, previous studies have investigated the oxidation behavior of Ni-based superalloy powders under different conditions. Xu et al. [22] investigated the oxidation behavior of Ni-based superalloy powders under different temperature and found that the oxygen content of the powders is low at room temperature but increases significantly at 250°C, with the powder surface being completely covered by an oxidation layer mainly composed of Ni(OH)2, Cr2O3, and TiO2. Zhang et al. [16] found that the NiO/Ni(OH)2 layer in the range of 5.7 to 9.8 nm exists on the Ni-based superalloy powder surface with different powder compositions and particle sizes by X-ray photoelectron spectroscopy (XPS). Normally, XPS has been widely used to identify the composition of the oxide layer on the surface of alloy powders due to its sensitivity to elemental valence states, not only in Ni-based superalloys, but also in other alloy powders [23–25]. Nevertheless, the detailed analysis of the oxide layer in powder surface still cannot establish a relationship between the microstructure and composition. Especially in Ni-based superalloy powders with diversity of alloying elements, the structure of oxide layer on the powder surface should be characterized in detail by more advanced methods [17].
To understand the low temperature oxidation behavior caused by the friction heat between the powder particles during the screening in the air in the actual production process, we artificially utilized different concentrations of oxygen to accelerate the oxidation of the alloy powders at low temperatures. This approach helps to solve the problem of characterization difficulties induced by the low oxygen content. Then, combined with the detailed characterization of transmission electron microscopy (TEM) and three-dimensional atomic probe (3D-AP), the features of the oxide layer were discussed. Finally, the formation mechanism of the oxide in different states was discussed. This work provides valuable information for the effectively characterizing the oxide film on the FGH96 superalloy powders, and contributes to the understanding of the oxidation behavior of the Ni-based superalloy powders.
The raw FGH96 alloy powders were produced by the vacuum induction melting argon gas atomization. Fig. 1 displays the particle size distribution of the powder used in this work. In view of cost, all the alloy powders below 50 μm were screened only once, and no secondary screening was performed to exclude finer alloy powder particles. As the powder size strongly affect powder reactivity and oxidation, the powders with particle size around 50 μm were specially investigated in this study. To obtain alloy powders with various oxygen contents, the powders were heated in a tube furnace at different temperatures in air atmosphere. The powders with medium and high oxygen contents were oxidated at 150 and 200°C for 30 min, respectively. The oxidation parameter and the oxygen content of the powders are shown in Table 1. The raw powders and that after oxidation were stored under high vacuum to prevent from further oxidation by the air.
| Sample | Process | Oxygen content / ppm |
| Low oxygen | Raw | ~140 |
| Medium oxygen | 150°C-30 min | ~280 |
| High oxygen | 200°C-30 min | ~340 |
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The oxygen contents of the powder were analyzed by a O/N/H analyzer (TC 600). The morphology of the alloy powders with various oxygen contents was investigated by a Regulus8100 field emission scanning electron microscope (FESEM). The particle size distribution of the alloy powders was measured and statistically analyzed by Nano Measure software based on the SEM images. The cross-sections of the powder surface were observed by transmission electron microscopy (TEM, Talors F200X G2) equipped with a high angle annular dark field (HAADF) detector. To distinguish the nanoscale phases of the powder surface, scanning transmission electron microscopy (STEM) mode was utilized. Furthermore, selected area electron diffraction (SAED) and nano-beam electron diffraction (NED) with a super energy dispersive spectroscope (EDS) were used to analyze the crystal structure and element distribution of the oxide film above powder surface. The qualitative analysis of composition was carried out by combining the point, line, and mapping analyses of the EDS. Besides, three-dimensional atomic probe (3DAP，LEAP 5000 XR) was adopted to further clarify the element distribution of the oxide layer in atomic scale, and CAMECA comprehensive visual analysis software IVAS was employed for data processing.
The TEM and 3DAP samples were prepared by focused ion/electron dual beam system (FIB, Helios 5 CX). It should be noted that the carbon (C) coating was deposited over powder surface to protect the oxide film of alloy powders, and this can also avoid the interference of metal coating elements on the alloy elements in the oxide film during the analysis. When preparing 3DAP sample, a copper (Cu) layer was pre-coated on powder surface due to the similar evaporation field with powder matrix (30 V/nm for Cu and 35 V/nm for Ni).
SEM images of FGH96 powders with different oxygen contents are displayed in Fig. 2. The powders with comparable sizes (~50 μm) are shown in Fig. 2(a), (c), and (e) respectively. The corresponding magnified images of the powder surface (Fig. 2(b), (d), and (f)) reveal the presence of three distinct regions, namely big precipitate, fine particulates, and matrix, which were marked in circle, rectangle and pentagon, respectively. The size of the big precipitates for all alloy powders with various oxygen contents ranges from 200 to 400 nm, while the size of the fine particulates is between 30 and 50 nm. In contrast, the matrix exhibits no apparent characteristics except for a clean surface. To clarify this phenomenon of the powder surface, more powders with different particle sizes (5–40 μm) were also characterized (Figs. S1 and S2), and the results show the same morphology.
The TEM images in Fig. 3 present the detailed surface microstructure of the alloy powders with various oxygen contents. Firstly, the thin film, as marked by white lines, is verified to be an oxide layer by the high oxygen concentration in EDS results (P1, P4, and P7). Due to the large surface area of the oxide layer in the TEM species, it can be concluded that the thin oxide layer is the “matrix” observed in SEM. Secondly, the big precipitates display a size range of 200–400 nm, which is consistent with that observed in SEM. From the EDS results in Table 2, a high concentration of Ti, Nb, and Mo elements is enriched in the precipitates (P2, P5, and P8) compared with that in powder matrix (P3, P6, and P9), suggesting that it is MC [12]. It should be noted that the concentration of the C element is seriously influenced by the C coating during FIB, but the preliminary qualitative analysis of the typical phase is not affected due to its comparable concentration of alloy elements with powder matrix. Thirdly, small particles with size of about 30 nm, corresponding to the fine particulates observed in SEM, are also found on the powder surface within medium oxygen contents, as shown in Fig. S3. These small particles are not observed in other samples due to their small size, which is difficult to be accurately positioned in FIB. The EDS results in Table S1 confirm that the element concentrations are the same as those in big precipitates, indicating that the small particulate is also MC. Therefore, the subsequent detailed analysis will be summarized into two categories: oxide layer over powder matrix and MC phase.
| Samples | Point | Element concentration / at% | ||||||||||
| C | O | Al | Ti | Cr | Co | Ni | Zr | Nb | Mo | W | ||
| Low oxygen | P1 | 40.10 | 29.44 | 2.75 | 4.65 | 6.20 | 2.62 | 9.87 | 1.30 | 0.17 | 1.72 | 1.18 |
| P2 | 32.65 | 4.65 | 2.22 | 26.37 | 6.10 | 1.67 | 4.19 | 0.96 | 8.11 | 10.72 | 2.35 | |
| P3 | 6.22 | 0 | 1.89 | 2.66 | 19.96 | 12.71 | 51.45 | 0.25 | 1.49 | 0.99 | 2.39 | |
| Medium oxygen | P4 | 33.77 | 27.34 | 2.83 | 4.49 | 5.18 | 4.08 | 19.85 | 1.43 | 0.38 | 0.43 | 0.23 |
| P5 | 34.13 | 4.39 | 0 | 33.17 | 2.55 | 1.00 | 3.26 | 1.62 | 8.30 | 7.12 | 4.45 | |
| P6 | 8.47 | 0 | 3.22 | 5.16 | 14.67 | 10.68 | 53.10 | 0.68 | 0.29 | 2.45 | 1.28 | |
| High oxygen | P7 | 29.80 | 27.62 | 1.91 | 2.10 | 7.01 | 4.71 | 23.40 | 0.83 | 0.48 | 1.83 | 0.31 |
| P8 | 41.41 | 0 | 1.67 | 27.52 | 3.31 | 2.83 | 5.69 | 2.17 | 6.40 | 5.84 | 3.15 | |
| P9 | 10.61 | 0 | 4.21 | 4.10 | 15.88 | 11.05 | 49.29 | 0.58 | 1.26 | 1.92 | 1.11 | |
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HRTEM was employed to characterize the features of the oxide layer over the powder matrix within various oxygen contents. Fig. 4(a)–(c) shows the oxide layer thicknesses of low, medium, and high oxygen content powders which are approximately 9, 14, and 30 nm, respectively. The thickness of the oxide layer was measured based on the contrast among the C coating and powder matrix, which can provide a relatively clear interface between the oxide layer and the powder matrix or C coating. Due to errors in manual measurements, 30 sets of data were taken to obtain the average thicknesses of the oxide layer. The result is shown in Fig. S4. Besides, the structure characteristics of the oxide layer were revealed by the NED patterns and fast Fourier transform (FFT) images, as shown in Fig. 4(a1)–(c1). Firstly, the diffraction patterns of the oxide layer and the powder matrix region display the typical face-centered cubic (FCC) structure under the zone axis of z = [011], which were confirmed as the γ matrix by the diffraction spot calibration. Secondly, an amorphous ring appears beside the diffraction spots of (111) γ crystal planes, as pointed by the white arrow in Fig. 4(a1) and 4(b1) indicating that the structure of the oxide layer should be amorphous. The diffraction pattern of the oxide layer of the alloy powder with high oxygen content only represents an amorphous ring (Fig. 4(c1)), verifying the amorphous structure of the oxide layer. It is necessary to note that the amorphous structure of oxide layer in alloy powder surface has been discovered in 316 L stainless steel [26−27].
Additionally, EDS mapping (Figs. S5–S7) presents the existence of the oxide layer due to the segregation of the O elements over the matrix. The line analysis of EDS (Fig. 4(d)) revealed the distribution of the alloy elements in the oxide layer of the powder with low oxygen content, where the alloying elements distributed layer by layer from the outer surface to the inner layer (Ni → Co → Cr → Al/Ti) in the range of oxygen rich region. The distribution of the alloy elements in the oxide layer in the other oxygen content powders remains the same (Fig. S8). It has been previously reported that multiple oxide layers are formed on the surface of alloys in succession due to the different oxidation rates of different elements [28−29]. Therefore, the amorphous oxide layer can be regarded as being composed of laminated oxides layers wrapped in the surface layer of the powder matrix. Moreover, the other oxide species of the oxide layer on the surface of FGH96 powders are consistent with those reported in Ref. [16].
To confirm the laminated oxides layers over powder matrix, 3DAP was used to more accurately analyze the delamination of alloying elements within the oxide layer at atomic scale. Fig. 5(a) presents the 3DAP image of the oxide layer of FGH96 powder with low oxygen content, and the atom concentration diagram (Fig. 5(b)) shows that the alloy elements of the powder matrix remain uniform, suggesting no γ′ precipitated in the powder matrix during gas atomization. This is consistent with the diffraction pattern analysis of the powder matrix investigated by TEM. In addition, Fig. 5(c) shows the reconstruction image of the 3DAP sample to clearly reveal the oxide layer. The concentration profile of elements in the oxide layer is shown in Fig. 5(d), and it can be seen that the alloying elements have a sequential concentration fluctuation between the rich regions of oxygen. However, it is difficult to determine the starting position of the fluctuation of element content. Therefore, the interval between the X-axis coordinate of the first peak and the X-axis coordinate of the decay to the first trough value was defined as the enrichment area of each element, although the range is smaller than the actual range. The corresponding data processing results are shown in the Fig. 5(e), and it can be found that there is a phenomenon of laminated distribution of elements in the range of oxygen atom concentration, which is followed by Ni rich area, Co rich area, Cr rich area, and Al rich area. The occurrence order of enriched alloying elements is consistent with EDS results observed under TEM. Herein, the optimized model of the oxide layer above the powder matrix is shown in Fig. 5(f), and the outermost of the laminated structure should be NiO layer, which is in accordance with Ref. [16].
Fig. 6 displays the TEM images of the carbides phase (as marked by the yellow rectangular) on the surface of FGH96 alloy powders with different oxygen contents. As shown in Fig. 6(a), the HRTEM image of region 1 shows the crystal structure of carbide in the alloy powder with the low oxygen content, and the SAED pattern revels the typical FCC structure with the zone axis of [001]. Through the calibration of the crystal planes, MC phase was confirmed by the (200) plane with a d-spacing of 0.213 nm. Besides, SAED pattern also presents an amorphous ring along with the diffraction spots of (200)MC, as marked by the white circle. Since the d-spacing of (200)MC (0.213 nm) is similar to that of (111)γ (0.208 nm), the amorphous ring appears in the SAED pattern of MC should be the same as the oxide layer, that is, an amorphous layer exists in region 1. Furthermore, the SAED patterns of the carbides in the other alloy powders with medium and high oxygen contents display the same result after the confirmation of MC phase (SAED images of Fig. 6(b) and (c)). Specially, Fig. 6(c) clearly reveals the amorphous layer (marked by white lines) over the MC phase according to the contrast of the TEM image.
EDS mapping was adopted to characterize the composition of the amorphous layer above the MC phase. As shown in Fig. S9, the EDS mapping images of the MC region of the powder surface with low oxygen content show the high concentration of Ti, Mo, Nb, and C elements, which reveals the position of MC phase. Fig. 6(d) displays the elements concentration of the region close to MC phase, and the result reveals the powder surface containing MC phase was divided into three regions from the outside to the inside. The inner region concentrates high percent of Ni, representing powder matrix (PM). The high concentration of Ti, Mo, Nb, and C elements suggests the MC phase in the middle region, while the high concentration of element O in the outer part implies the oxide layer. Combining with the EDS mapping result to analyze, it is interesting to find that only the Ti enriches in the oxide layer region. Therefore, it can be inferred that the amorphous layer observed over MC phase in TEM should be Ti rich oxide layer. Furthermore, the EDS results of the other two powders with medium and high oxygen contents also reveal the segregation of Ti in the oxide layer (Figs. S10–S12), especially in Fig. S11 of the powder with high oxygen content. Considering the SEM image of the precipitates, it can be concluded that the oxide layer over MC phase should be amorphous Ti rich oxide particle, as presented in the model diagram in Fig. 6(e). Moreover, the thickness of the oxide particles over the MC phase ranges from around 10 to 30 nm with the increase of the oxygen content of the alloy powder, which is consistent with the oxide layer over powder matrix. The different segregation of the alloy elements in the oxide layer may be attributed to the different precipitation behaviors of the elements during the gas atomization, which may affect the evolution of the oxygen in the subsequent process [15].
The oxidation of the alloys is affected by many factors, and the metals in the alloy will have different affinities for oxygen, leading to the different oxidation behavior of alloy elements [30]. As a consequence, the different compositions of the oxide layer covering the matrix and MC phase may have the different oxidation mechanisms.
As for the oxide layer covering the matrix, the varying alloy elements suggest the oxidation behavior is complex. Normally, Gibbs free energy (∆G) versus composition diagrams is used to thermodynamically determine the possibility of the composition of the oxide layer [30]. Firstly, as shown in Fig. 7, the ∆G for the reaction 2Ni+O2=2NiO is negative, suggesting that the powder matrix has the solubility and diffusivity for oxygen that is sufficient to establish the desired dissolved O activity at the oxide layer [30]. In addition, this can be taken as evidence of an internal diffusion of oxygen. Secondly, it can be seen from Fig. 7, the ∆G for forming oxides (CoO, Cr2O3, TiO2, and Al2O3) are more negative than that for forming NiO (per mole), indicating that the formation of these oxides is easier than that of NiO. However, the amorphous structure of the oxide layer (Fig. 4) indicates that the oxides do not sufficiently nucleate and grow due to the rapid solidification when the alloy powder suffers atomization freezes [27]. Thirdly, the stabilities of the oxides can be seen directly from the Ellingham diagram, and the lower position in the diagram corresponds to the more stable oxide. The equilibrium oxygen partial pressure of the oxide (PoxideO2) could be calculated by the equation: PoxideO2=exp[ΔGoxide/(RT)] (where ΔGoxide, R, and T are the Gibbs free energy for oxides, gas constant, and temperature, respectively), which implies the lower the Gibbs free energy of the oxide, the lower the partial pressure of oxygen required to form the oxide. However, the oxide would decompose when the oxygen partial pressure falls below the equilibrium pressure required for the formation of the oxide [31], and the formation of the layered structure is attributed to the decreasing equilibrium pressure as oxygen diffuses from the powder surface into the matrix [32]. The layered structure of the oxide layer for the alloy powders examined by electron energy loss spectroscopy is also found in 316 stainless steel powders [27].
As for the oxidation of the oxide particles over MC phase, the oxide particles with more intense oxidation appear on the powder surface with high oxygen content. Fig. 8(a) displays the macro morphology of the oxide particles, as marked in black rectangle. The enlarged TEM image (Fig. 8(b)) shows the strip morphology of the particle, and the size of it is around 200 nm, which is similar to MC phase discussed earlier. However, EDS result in Table 3 reveals that the enriched oxygen atoms in strip particles reach 28.2% with 11.6% Ti atoms, suggesting that there may be TiO2. Fig. 8(c) exhibits the TEM image in dark field of the TiO2 particle, and it can be inferred that the TiO2 particle possesses a second phase according to the bright contrast of the image. Nevertheless, the SAED pattern (Fig. 8(d)) of the TiO2 particle displays an amorphous ring, implying it possesses an amorphous structure.
| Sample | Point | Element concentration / at% | ||||||||||
| C | O | Al | Ti | Cr | Co | Ni | Zr | Nb | Mo | W | ||
| High oxygen | P10 | 44.48 | 28.2 | 0.13 | 11.16 | 1.48 | 0.95 | 2.14 | 0.33 | 3.85 | 5.04 | 2.24 |
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The EDS mapping (Fig. 8(e)) and line analysis (Fig. 8(f)) reveal the same result as point analysis of the TiO2 particle, while the segregation of a small amount of Mo, Nb, and C elements appears in the inner part of the particle. It can be found that these elements are also the main elements that constitute the MC phase (Fig. 6(d)), implying a transformation between the MC phase and the TiO2 particle. Besides, as shown in Table 3, the atomic concentration of carbide-forming elements such as Ti, Nb, and Mo elements in the oxide particle (P10 in Fig. 8) was found to be lower than that in the MC phase (P8 in Fig. 3(e)), and it can be seen from Fig. 8(f) that the interval of the oxide particle reaches 60 nm, which is much wider than that of the oxide layer over the MC phase (30 nm). Consequently, it can be inferred that Ti atoms in the MC particles are consumed by O atoms with the exacerbation of the oxidation, leading to the narrowing in the size of MC phase. Therefore, the TiO2 particle could be categorized as the peroxide of the MC phase during oxidation due to the same composition.
The oxidation mechanism of the TiO2 particle is strongly associated with the MC phase [33]. Zhao et al. [34] have investigated the oxidation behavior of TiC and Ni in TiC/Ni composites. They found that TiC decomposes, leading to the outward diffusion of Ti atoms. Consequently, TiO2 islands form over the TiC. Besides, Cai et al. [35] testified the final oxidation product of the (Ti,Nb)C at high temperatures should be TiO2 and CO and CO2. These works all suggest the possibility of the oxidation behavior of TiC in TiC/Ni composites. In this work, thermodynamic analysis results (Fig. 9) show that the ∆G of reaction between Ti and O is lower than that between Ti and C, implying the Ti element is more likely react with O other than C. Moreover, the ∆G of the reaction between TiC and O2, whether the oxidation product is CO or CO2, is lower than that of the formation of TiC between Ti and C. This could be the evidence for the presence of TiO2 above MC phase. The EDS analysis in Fig. 3 also reveals the segregation of the O in MC phase. Combined with Fig. 8 to analyze, the dramatic reduction in the size of the MC in amorphous oxide particle implies the TiC may completely decompose to form TiO2 when the oxidation process is prolonged.
Based on the above analysis of the oxide layer of powder surface, it can be concluded that the oxidation behavior of the alloy powder was influenced by the MC phase precipitated over powder surface during the gas atomization [36]. The oxidation behavior FGH96 alloy powder should be divided into two types: the oxidation of the γ matrix and the MC phase (Fig. 10(a). Two different oxidation types result in different morphologies of oxides on the powder surface. As the oxygen content of the alloy powder increases, the thickness of the surface oxide film, whether it is the oxide layer above the powder matrix or the oxide particles above the MC phase, also increases correspondingly (Fig. 10(b)).
Moreover, the reconstruction image provides a clear visualization of the structure and composition of the oxide layer and suggests possible interface scenarios during the subsequent consolidation process. These scenarios include: (1) the contacting between the oxide layer over powder matrix of powder 1 and that of powder 2; (2) the contacting between oxide particle over MC of powder 1 and that of powder 2, and (3) the contacting between the oxide particle over MC of powder 1 and the oxide layer over powder matrix of powder 2. The different alloying elements present in the oxide layer and oxide particle may result in diverse genetic evolution of the oxygen on the powder surface, which will be analyzed in detail in future studies.
The microstructure of the oxide film on the surface of FGH96 alloy powders with varying oxygen content was investigated using advanced techniques such as FIB cutting, HRTEM, and 3DAP. The findings can be concluded as follows.
(1) The oxide film of the FGH96 alloy powders could be categorized into two parts: the oxide layer over the γ matrix and the oxide particle situated above the MC phase, both of which exhibit an amorphous structure.
(2) The alloying elements within the amorphous oxide layer exhibit a laminated distribution, with Ni, Co, Cr, Ti/Al from the surface to the γ matrix. This is because of the decreasing oxygen equilibrium pressure as oxygen diffuses into the matrix.
(3) The oxide particles located above the MC phase primarily contain a high concentration of Ti elements, and thermodynamic calculations confirm these particles to be peroxide products of the MC phase.
(4) The distribution of oxides in the surface oxide layer of powders with varying oxygen contents is similar, while the thickness of the amorphous oxide film on the powder surface increases from 10 to 30 nm with an increase in oxygen content in FGH96 alloy powders.
This work is financially supported by the National Key R&D Program of China (No. 2021YFB3704000), the National Natural Science Foundation of China (Nos. 52074032, 51974029, 52071013, and 52130407), the Beijing Natural Science Foundation (No. 2232084), the Guangdong Basic and Applied Basic Research Foundation (No. 2021B1515120033), the 111 Project (No. B170003), and the Basic and Applied Basic Research Fund of Guangdong Province, China (No. BK20BE015).
Xuanhui Qu is an editorial board member for this journal and was not involved in the editorial review or the decision to publish this article. The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.
The online version contains supplementary material available at https://doi.org/10.1007/s12613-024-2823-x.
*These authors contribute equally to this work.
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| 1. | Yang Liu, Yufeng Liu, Lin Zhang, et al. Investigating the Dissolution and Evolution Behavior of Oxides in FGH96 Superalloys. Metallurgical and Materials Transactions A, 2025, 56(1): 140. DOI:10.1007/s11661-024-07621-3 |
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Figures(10) / Tables(3)
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| Sample | Process | Oxygen content / ppm |
| Low oxygen | Raw | ~140 |
| Medium oxygen | 150°C-30 min | ~280 |
| High oxygen | 200°C-30 min | ~340 |
DownLoad:
CSV
| Samples | Point | Element concentration / at% | ||||||||||
| C | O | Al | Ti | Cr | Co | Ni | Zr | Nb | Mo | W | ||
| Low oxygen | P1 | 40.10 | 29.44 | 2.75 | 4.65 | 6.20 | 2.62 | 9.87 | 1.30 | 0.17 | 1.72 | 1.18 |
| P2 | 32.65 | 4.65 | 2.22 | 26.37 | 6.10 | 1.67 | 4.19 | 0.96 | 8.11 | 10.72 | 2.35 | |
| P3 | 6.22 | 0 | 1.89 | 2.66 | 19.96 | 12.71 | 51.45 | 0.25 | 1.49 | 0.99 | 2.39 | |
| Medium oxygen | P4 | 33.77 | 27.34 | 2.83 | 4.49 | 5.18 | 4.08 | 19.85 | 1.43 | 0.38 | 0.43 | 0.23 |
| P5 | 34.13 | 4.39 | 0 | 33.17 | 2.55 | 1.00 | 3.26 | 1.62 | 8.30 | 7.12 | 4.45 | |
| P6 | 8.47 | 0 | 3.22 | 5.16 | 14.67 | 10.68 | 53.10 | 0.68 | 0.29 | 2.45 | 1.28 | |
| High oxygen | P7 | 29.80 | 27.62 | 1.91 | 2.10 | 7.01 | 4.71 | 23.40 | 0.83 | 0.48 | 1.83 | 0.31 |
| P8 | 41.41 | 0 | 1.67 | 27.52 | 3.31 | 2.83 | 5.69 | 2.17 | 6.40 | 5.84 | 3.15 | |
| P9 | 10.61 | 0 | 4.21 | 4.10 | 15.88 | 11.05 | 49.29 | 0.58 | 1.26 | 1.92 | 1.11 | |
DownLoad:
CSV
| Sample | Point | Element concentration / at% | ||||||||||
| C | O | Al | Ti | Cr | Co | Ni | Zr | Nb | Mo | W | ||
| High oxygen | P10 | 44.48 | 28.2 | 0.13 | 11.16 | 1.48 | 0.95 | 2.14 | 0.33 | 3.85 | 5.04 | 2.24 |
DownLoad:
CSV
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