The formation and evolution of oxide particles in oxide-dispersion-strengthened ferritic steels during processing

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Acta Materialia 61 (2013) 2219–2235 www.elsevier.com/locate/actamat

The formation and evolution of oxide particles in oxide-dispersion-strengthened ferritic steels during processing Ceri A. Williams a,⇑, Paulina Unifantowicz b, Nadine Baluc b, George D.W. Smith a, Emmanuelle A. Marquis a,c b

a Department of Materials, University of Oxford, Oxford OX1 3PH, UK Ecole Polytechnique Federale de Lausanne (EPFL), Centre de Recherches en Physique Plasmas, Association Euratom-Confederation Suisse, 5232 Villigen PSI, Switzerland c Department of Materials Science and Engineering, University of Michigan, Ann Arbor, MI 48109-2136, USA

Received 26 July 2012; received in revised form 21 December 2012; accepted 23 December 2012 Available online 29 January 2013

Abstract The fabrication of oxide-dispersion-strengthened (ODS) steels is a multi-stage process involving powder ball milling, degassing and consolidation by hot isostatic pressing. Y is introduced by mechanical alloying (MA) with either Y2O3 or Fe2Y so a high density of Y–Ti–O-based oxide nanoparticles is formed. The evolution of 2 nm oxide nanoparticles and larger >5 nm grain boundary oxides has been studied at each step of the processing. The nanoparticle dispersions produced in material MA with Fe2Y were comparable to those produced by alloying with Y2O3. Hence the majority of oxygen which forms the nanoparticles must be incorporated from the atmosphere during MA, rather than being introduced via the alloying additions. During mechanical alloying, a high density of subnanometer particles are formed (2.5 ± 0.5  1024 m3). The oxygen content of the nanoparticles decreases slightly on annealing, and then the composition of the nanoparticles remains constant throughout subsequent processing stages. The nanoparticle size increases during processing up to 2 nm radius, while the number density decreases to 4 ± 0.5  1023 m3. Grain boundary oxides (>5 nm) have a Ti–Cr–O-rich shell, and contain no Y before consolidation, but have similar core composition to the matrix nanoparticles after consolidation. The formation of the larger grain boundary oxides is shown to take place during the degassing and consolidation stages, and this occurs at the expense of the nanoparticles in the matrix. This work provides a mechanistic understanding of the importance of controlling the oxygen content in the powder during MA, and the resulting impact on the formation of the ODS microstructure. Ó 2013 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Oxide dispersion strengthening; Ferritc steel; Grain boundary; Atom probe

1. Introduction Oxide-dispersion-strengthened (ODS) steels are promising candidate materials for structural applications in nuclear fusion reactors and fourth generation fission reactors due to their high temperature stability and strength, and resistance to radiation damage. The manufacturing of ODS steels is an energy-intensive process commonly involving mechanical alloying (MA) of the base alloy with ⇑ Corresponding author. Tel.: +44 1865 273634.

E-mail address: [email protected] (C.A. Williams).

Y2O3, then consolidation by hot isostatic pressing (HIP) or extrusion. This produces a material with a small grain size and a fine dispersion of nanoscale oxide particles that has been shown to greatly reduce the detrimental effect of irradiation [1–3]. A comprehensive review of the development of ODS steels by Odette et al. can be found in Ref. [4]. The distribution of oxide particles can vary substantially depending on the composition of the alloy. Aluminumcontaining alloys such as PM2000 typically show a distribution of Y–Al–O-based particles 10–50 nm in diameter [5]. Ferritic/martensitic alloys such as ODS-Eurofer and the 12-YW alloy often have a dispersion of Y–O-based

1359-6454/$36.00 Ó 2013 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.actamat.2012.12.042

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particles 5–20 nm in diameter [6,7]. It has been shown that Ti additions play an important role in refining the particle dispersion to 2 nm nanoparticles rich in Ti, Y and O [8,9] but the exact mechanism of this effect is unclear [10]. There is considerable controversy over the structure of the oxide particles. A wide variety of particle structures has been reported in various alloys. Transmission electron microscopy (TEM) analyses showed that some particles larger than 15 nm were consistent with the Y2Ti2O7 phase in MA957 [11]. The same structure was observed in a Ticontaining ODS-Eurofer alloy for particles larger than 5 nm [12]. Small-angle neutron scattering (SANS) data obtained from the 14YWT alloy [8] and TEM analysis on 18% Cr containing steel suggested a mixture of Y2TiO5 and Y2Ti2O7 [13–16]. Cayron et al. also suggested the presence of a new structure, ”pyro-ortho”, consisting of a complex mixture of Y2O3 and pyrochlore phases [17]. De Castro et al. investigated a model alloy containing no Ti, and observed particles having structures consistent with Y2O3, YCrO3 and Cr2O3, although 60% of particles could not be identified as any known oxide structure [18]. High resolution TEM analysis of nanoparticles >2 nm radius in the 14-YWT alloy has been equally ambiguous. Sakasegawa et al. suggested that nanoparticles in the MA 957 material were non-stoichiometric [11]. Wu et al., however, observed near stoichiometric Y2Ti2O7 nanoparticles [19] and Ribis and de Carlan also observed particles with structures corresponding to Y2Ti2O7 after annealing at 1300 °C for 1 h [20]. Hirata et al. claimed that the nanoparticles had a rock-salt-type structure [21], whereas Brandes et al. indicated that the nanoparticles were in fact amorphous [22]. Energy-dispersive X-ray analysis of particles in a Fe–16 Cr–0.1 Ti ODS steel by Kishimoto et al. indicated a Y:Ti ratio of 2:1, which corresponds to the Y2TiO5 oxide [15]. Electron energy loss spectroscopy measurements have been reported to show that the particles are consistent with Y2Ti2O7 [16]. For particles with radii <2 nm, Miller et al. measured a Y:Ti ratio of considerably less than 1 (ranging from 0.25 to 0.5) by atom probe tomography (APT), with an overall metal to oxygen ratio (M:O) of 1:1 [23–25], which would correspond to an oxide with TiO type structure. This was in agreement with the findings of Hirata et al. [21]. There is considerable variation in the composition of oxide nanoparticles measured by APT, even when the same material is under investigation [23–25]. There is, however, some agreement in the M:O ratio, and a 1:1 ratio is commonly observed. Even in the ODS-Eurofer material, an M:O ratio of 1:1 was measured by APT [26], but in this case a TiO type structure was thought to be very unlikely as the alloy composition contained no Ti. A common interpretation of the substoichiometeric oxygen content measured by APT is that not all the oxygen is detected during the APT data acquisition [19], but a recent study has shown that this may not be the case [10]. Instead it appears that effects such as multiple ion evaporation events can in some cases lead to a Y deficiency.

This is discussed further in Section 4.1. APT results suggest that many ODS particles have a core–shell-type structure [26,27], which will also affect the apparent composition measurements. The shell structure has also been observed by TEM [28–30]. One theory for the formation of the oxide particles is that the shell acts as a wetting layer for the nucleation of the particles [26], and it was proposed by Sakasegawa et al. that the formation mechanism is strongly dependent on the diffusion of Y [28]. In order to optimize alloy performance and refine processing routes, it is vital to establish the mechanism of formation of the oxide nanoparticles, and their subsequent development, and to understand the role of different alloying elements on their behavior. SANS data by Alinger et al. on 14-YWT (Fe–14 Cr–2 W–0.3 Ti) suggested complete dissolution of Y and O during the MA process and precipitation of the oxide nanoparticles during HIP [10]. However, Brocq et al. observed partitioning of Ti–Y–O in powder directly after MA in an alloy with similar base composition [31]. Brocq et al. also showed that the particles continued to nucleate and grow during relatively low temperature annealing (<800 °C) for short times (5 min), and suggested that these particles formed the basis for the distribution seen in consolidated materials. The variety of alloy compositions and processing conditions in the previously published studies makes it difficult to determine whether the variability in oxide phase formation is due to differences in composition and processing, or is inherent to the nature and stability of the oxide nanoparticles. Characterizing the microstructure of well controlled alloy systems at various stages during the processing will assist in evaluating the mechanism of formation of the nanoparticles, and will permit better control of the processing conditions. As the manufacturing is a multi-stage process, it is important to understand what influence each stage of the processing has on the final microstructure. It was suggested by Brocq et al. that MA using Fe–Y compounds to introduce Y could generate an ODS microstructure that is very similar to that generated by MA with Y2O3 [31]. Oksiuta et al. used Fe2Y during mechanically alloying, and reported that less oxygen was incorporated into the alloy [32]. The authors suggested that the reduction in O may be responsible for the generation of oxide nanoparticles slightly larger in diameter (6 nm) than the 2 nm particles seen when MA with Y2O3 [32]. In this study, the microstructure of a Fe–14 Cr–2 W–0.3 Ti ODS alloy MA with 0.3 wt.% Y2O3 (designated EPFLE-Y2O3) was studied using both APT and TEM. The powder was characterized directly after MA, at several intermediate stages of processing, and after a post-consolidation heat treatment to elucidate the evolution of the oxide nanoparticles and larger grain boundary oxides. Another alloy with near-identical processing conditions but MA with 0.5 wt.% Fe2Y (designated EPFL-E-Fe2Y) was also studied, to compare with EPFL-E-Y2O3 alloy at various stages of processing. This allowed the way in which Y and O are incorporated into the matrix to be investigated.

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2. Experimental details Elemental powders of Fe, Cr, W and Ti were mixed with either Y2O3 or Fe2Y and milled in a planetary ball mill at 300 rpm for 50 h in an H2 atmosphere to produce an alloy with a nominal base composition of Fe–14 Cr–2 W–0.3 Ti and either 0.3 wt.% Y2O3 or 0.5 wt.% Fe2Y. The powder was degassed at 800 °C for 2 h, and consolidation was achieved by hot isostatic pressing at 2 MPa, 1150 °C for 4 h. Focused ion beam milling (FIB) with a Zeiss NVision 40 with Kleindiek micromanipulator was used to attach sections of powder particles to silicon posts, and annular milling was used to reduce these sections to needle-shaped specimens suitable for APT [33]. Sections of powder were also attached to copper TEM half-grids using FIB, and thinned to <100 nm to make specimens suitable for TEM. 40 kV Ga ions were used to machine the specimens, with the final milling stage carried out at 2 kV to avoid excessive implantation of Ga. For the consolidated material, standard electropolishing techniques as described in Ref. [26] were used to prepare specimens suitable for TEM and APT. TEM observations were made using a Philips (FEI) CM20 microscope or a Hitachi HD2700 operated at 200 kV, and electron microprobe analysis (EPMA) was carried out using a JEOL JXA-8800 instrument operating at 20 kV using a current of 50 nA for 100 s count time. The APT analysis was carried out using an Imago LEAP 3000 HR operating in laser pulsing mode. An APT specimen base temperature of 30 K was used with a laser energy of 0.3–0.4 nJ and a laser spot size of 8 lm, at a repetition rate of 200 kHz. The interpretation of the APT results relies on the analysis of clustering across multiple length scales. Different data analysis methods were employed to assess the size, composition and number density of the particles, depending on their size. For large particles, where the core of the oxide feature is >5 nm, the particles were isolated by generating an iso-concentration surface corresponding to 1% Ti (including Ti contained in molecular ion species), and the composition of the particles and interface chemistry were analyzed using the proxigram method [34]. These larger oxide particles were then removed from the data reconstruction to investigate clustering on a finer scale. Smaller nanoparticles often did not contain sufficient numbers of atoms to use the proxigram method reliably, therefore clustering was identified using the maximum separation cluster selection algorithm [35]. Due to the mechanical alloying process, the supersaturation of solutes was much higher than in classical systems, and there was substantial variation in the nearest neighbor distances of solutes depending on the extent to which the material had been heat treated after ball milling. This affected the choice of cluster selection parameters required to identify the clusters. A consistent method to define the cluster selection parameters was required for every three-dimensional

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(3-D) reconstruction, to ensure that clusters were accurately defined in each case. An Nmin value of 5 was used and dmax values ranging from 0.7 to 1.3 nm were chosen based on a comparison between the APT data, and data that were randomized by mass for each of the x–y–z coordinates [36,37]. For the purposes of this study, the size, composition, number density and volume fraction of the particles (as defined by the maximum separation method) were required. Due to trajectory aberrations associated with the lower evaporation field of the nanoparticles, the density of atoms detected in the regions associated with the nanoparticles increased by approximately 4, with the extra atoms originating from the matrix [38]. As in Ref. [26], the contribution of the matrix was assessed by examining the level of Fe in the clusters. However, the exact quantity of Fe in the clusters was not addressed in this work. Instead, the Fe level in the clusters was artificially set to zero, and the number of Fe atoms used to estimate the expected quantity of other matrix elements (Cr, W, etc.) introduced by trajectory aberrations. This matrix contribution was removed and the resulting cluster compositions are quoted as “matrix corrected”. The nanoparticle sizes were compared by calculating the average radius of gyration (Rg) as described in Ref. [34]. This measurement does not reflect the size of the nanoparticles measured with other techniques, but is reliable in detecting relative changes in nanoparticle size within APT data. Due to this uncertainty, the volume fraction occupied by the nanoclusters was calculated from the estimated number of cluster atoms after the “matrix correction” step. Segregation at grain boundaries was quantified using a one-dimensional (1-D) cylindrical concentration profile. The concentration was calculated at 0.5 nm intervals along a cylinder of 10–20 nm radius orientated with the length of the cylinder perpendicular to the grain boundary plane. By quantifying the amount of segregation using the Gibbsian interfacial excess [39], it is possible to compare segregation at each boundary directly. The Gibbsian interfacial excess is calculated using Ci ¼ nAix , where the area A of the interface can be approximated to the diameter of the cylinder used for analysis (as the area of cylinder is small, curvature of the interface can be assumed negligible). The total number of excess solute atoms nix is found from the integral of the composition profile across the grain boundary region above the matrix level, as described in Refs. [40,41]. The width of the interface is defined by the full width of the solute peak produced by the most strongly segregated element. The error in this value has been estimated as the standard error of at least three separate measurements at different points along the boundary. 3. Results The nominal compositions of the EPFL-E-Y2O3 and EPFL-E-Fe2Y materials and the compositions of both alloys as measured by EPMA are given in Table 1. There

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Table 1 Nominal composition of the consolidated EPFL-E-Y2O3 and EPFL-E-Fe2Y materials with corresponding EPMA analysis. All values are converted to at.%. Element

Nominal Y2O3

Y2O3 EPMA

Nominal Fe2Y

Fe2Y EPMA

Fe Cr W Y Ti O C N

83.6 15.1 0.6 0.15 0.4 0.21 – –

84.3 ± 0.2 14.73 ± 0.2 0.42 ± 0.1 0.13 ± 0.0 0.42 ± 0.0 – – –

82.9 15.0 0.6 0.14 0.4 – – –

84.9 ± 0.2 14.63 ± 0.2 0.44 ± 0.1 0.1 ± 0.1 0.37 ± 0.0 – – –

is good agreement between the nominal composition and the composition measured by EPMA in each case. There is therefore no significant difference in the overall solute concentration between the two materials. The oxide particle development was investigated by characterizing the microstructure at the following stages during the processing route for the EPFL-E-Y2O3 alloy. Comparative results will be presented for the EPFL-EFe2Y at stages 1 and 4 to show the effect of MA with the Fe2Y intermetallic to incorporate the Y. Stage 1: Powder, immediately after MA. Stage 2: Powder, after a short anneal at 800 °C (10 min). Stage 3: Powder, after degassing (annealing at 800 °C for 2 h). Stage 4: After degassing and consolidation by HIP at 1150 °C for 4 h. Stage 5: After degassing, consolidation and further annealing for 100 h at 1150 °C. An overview of the microstructural features at each stage of processing is first presented, followed by separate in-depth analysis of the matrix composition, oxide nanoparticles, grain boundaries and grain boundary oxides for both alloys. 3.1. Overview of microstructural development Fig. 1 shows the microstructural evolution in the powder after MA, in the powder after degassing, and in the consolidated material for the EPFL-E-Y2O3 alloy. The scanning electron microscopy (SEM) image of a sectioned powder flake after MA shows some porosity and small pieces of powder loosely fused together (Fig. 1a). TEM of the MA powder reveals a textured nanocrystalline structure, with the average distance between grain or subboundaries <20 nm (Fig. 1b). The nano-grained structure seen in a section of powder flake immediately after MA remains present in over 90% of the powder volume after degassing (Fig. 1c). A number of recrystallized grains are also observed in areas near the edge of the powder flakes after degassing (Fig. 1d). A dispersion of small oxide nanoclusters is observed within these grains, together with a small fraction of larger oxide particles 10–50 nm in diameter. The consolidated material is

fully recrystallized, and a mixture of very small grains <100 nm can be seen alongside grains of 1 lm (Fig. 1e). The average grain size is 80 ± 20 nm and these grains contain a bimodal distribution of oxide nanoclusters (Fig. 1f). There is no significant change in the microstructure after annealing for 100 h at 1150 °C. Very similar features are observed in the EPFL-E-Fe2Y alloy. Fig. 2a shows that in the powder after MA a textured nanocrystalline structure in observed, comparable to Fig. 1a. Again, the consolidated material is fully recrystallized, and shows a bimodal distribution of oxide nanoclusters (Fig. 2b). The EPFL-E-Fe2Y material has in excess of 80% of grains that are fully recovered, with few visible dislocations in those grains. There are only a very small number of grains recovered to the same extent in the EPFL-E-Y2O3 alloy, where the dislocation density remains much higher. The average grain size in the EPFL-E-Fe2Y material is also larger at 310 ± 50 nm. 3.2. APT analysis of the solute element distribution within grains 3.2.1. EPFL-E-Y2O3 Clustering of Y, Ti, O and Cr is evident at all processing stages, even in the powder immediately after MA. Fig. 3 highlights the differences in the extent of the clustering that occurs during processing. Initially, at stage 1, the clusters are at the <1 nm level, but grow considerably during the subsequent processing stages. The overall composition as measured by APT for each condition is given in Table 2, along with the matrix composition. Bulk composition measurements are calculated by counting the total number of atoms of each element in a complete APT analysis (typically 20–50 million ions), and the values are averaged over multiple datasets. The matrix compositions are calculated in a similar manner, but atoms defined as part of a cluster by the maximum separation algorithm are excluded. Substantial changes occur during processing for both the bulk composition and matrix composition measurements. The matrix composition shows a decline in the clusterforming elements from stage 1 to stage 5 of processing. Initially there is some O, Y and Ti in the matrix, but after stage 3 of processing this is reduced to a trace level. There is also a decline in the proportion of the cluster-forming elements in the bulk composition measurements. This is

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Fig. 1. SEM and bright-field TEM observations of the microstructural evolution of the EPFL-E-Y2O3 ODS steel: (a and b) stage 1, powder after MA; (c and d) stage 3, powder after degassing; (e and f) stage 4, consolidated material.

particularly evident in the oxygen concentration. A likely explanation for this is that a proportion of these elements form larger oxides which are often outside the APT field of view. The bulk composition measurements taken from the powder flakes are less consistent in terms of solute content than in the consolidated material, which is reflected in the larger errors on the Cr levels in particular. This is assumed to be due to slight variations in the composition of each powder flake following the MA process. The oxide nanoparticle compositions at each stage of processing are given in Table 3. In each condition, all the clusters contain Y, Ti, O and a large proportion of Cr, even after the matrix contribution has been removed. Significantly more oxygen is observed in the clusters immediately after MA. There is greater consistency between the nanoparticle compositions in the annealed powders and the consolidated materials. The proportion of Y in the nanoparticles increases by a factor of two after consolidation. The number density and volume fraction of the nanoparticles show substantial changes. The number density of clusters in the MA powder is initially very high, but

decreases slightly after annealing the powder for 10 min at 800 °C. As the annealing process is continued, the number density of clusters falls by a factor of 5, and after consolidation the number density reduces even further. The volume fraction occupied by the nanoclusters would be expected to increase during processing, as a decrease in Ti and O in the matrix is evident from Table 1. This trend is observed for the shortest annealing time, but for all subsequent stages of processing, Table 3 shows a continuous decrease in the measured volume fraction. The size distribution of the nanoparticles at each stage of processing shows a gradual coarsening from stage 1 to stage 3 as shown in Fig. 4. The increased temperature and pressure during consolidation means that there is a sharp increase in the mean radius of the nanoparticles between stages 3 and 4, but little coarsening is observed on subsequent annealing at 1100 °C for 100 h. 3.2.2. EPFL-E-Fe2Y There are remarkable similarities between the nanoparticles observed in the EPFL-E-Y2O3 and EPFL-E-Fe2Y

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contamination of the atmosphere during the MA stage. The number density and volume fraction (seen in Table 5) and the size distribution at stage 1 and stage 4 of processing (shown in Fig. 6) all indicate that the same mechanisms of nanoparticle evolution apparent in EPFL-E-Y2O3 are also occurring in EPFL-E-Fe2Y. 3.3. Grain boundary analysis

Fig. 2. Bright-field TEM image of the EPFL-E-Fe2Y material: (a) stage 1, powder after MA; (b) stage 4, consolidated material.

alloys, therefore only stages 1 and 4 were compared in detail. Fig. 5 shows that the same trends of Y, Ti and Cr clustering are observed as in Fig. 3. As with the EPMA measurements in Table 1, the APT measurements of bulk composition for the EPFL-E-Fe2Y alloy compare favorably with the composition of EPFL-E-Y2O3, as shown in Table 4. Again, the levels of Y, Ti, Cr and O are higher in the matrix in the powder directly after MA than in the consolidated material, comparable to the differences seen in EPFL-E-Y2O3. The composition of the nanoparticles in EPFL-E-Fe2Y is very similar to those in EPFL-E-Y2O3. The only significant difference is the slight increase in N content in EPFLE-Fe2Y powder after MA. This is likely to be due to

3.3.1. EPFL-E-Y2O3 Grain boundaries are readily observed in APT data if they exhibit solute segregation, or substantial atomic density differences creating a “lensing effect” due to the irregular surface topography during APT analysis. Despite analyzing more than 10 specimens, no solute segregation indicating the presence of grain boundaries was observed in the EPFL-E-Y2O3 powder directly after MA. Fig. 7a shows high density regions which correspond to the size of the nanograined structure visible in Fig. 1b. There is no significant solute segregation observable across these boundaries, as seen in the 1-D concentration profile in Fig. 7b. Grain boundaries exhibiting solute segregation were observed at all other stages of processing. The Gibbsian interfacial excess values from typical boundaries are shown in Table 6. The boundaries are usually enriched in Cr, W and C. An example of the extent of the solute segregation in the consolidated material is shown in Fig. 7c and d. Along with a uniform increase in Cr, W and C across the grain boundary plane, larger oxide particles (>5 nm) were observed in all the annealed powder conditions and consolidated material (stages 2–5). A high density region at the boundary corresponding to such a particle is indicated in Fig. 7c. In the powder directly after MA, no oxide particles larger than 2 nm diameter were observed either by TEM or by APT. It is possible that the complex nanocrystalline structure and high dislocation density seen in Fig. 1b obscure the larger oxides from view, but more likely that the formation of larger oxides largely occurs after the MA stage. While the compositions of the nanoparticles are comparable in each condition, those of the larger particles found at grain boundaries are seen to vary substantially depending on the stage of processing. Fig. 8 shows how the interface chemistry of particles >5 nm in diameter changes during processing. After annealing at 800 °C for 10 min, the largest observed particle size is 12 nm in diameter. The measured composition at the core of these 8–15 nm particles is 35% Cr and 55% O, which is consistent with a slightly oxygen-deficient Cr2O3 stoichiometry. Each particle has a shell of 2 nm thickness enriched in Ti, Cr and O. Where the powder was only annealed for 2 h at 800 °C, a number of particles 15 nm in diameter were observed. In these particles, the same core–shell structure is found; however, the Ti rich shell has grown to be 4–5 nm wide. In the case of the consolidated material, similar Ti–Cr–O particles are observed (5–15 nm diameter). For particles larger than 15 nm, the core–shell structure is also observed;

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Fig. 3. Five nanometer thick slice of an atom map showing the distribution of Cr, Y and Ti–O ions in the EPFL-E-Y2O3 material: (a) stage 1, powder after MA; (b) stage 4, consolidated material.

Table 2 APT measurements of bulk and matrix compositions for all stages (1–5) of processing in the EPFL-E-Y2O3 material. Bulk measurements are averaged from all atoms in at least three analysis volumes containing 30 M ions+. Matrix composition measurements are estimated by removing the particles using the maximum separation method, and only including areas away from grain boundaries that exhibit solute segregation. 800 °C 10 mins

After MA Bulk r Fe Cr W Y Ti O C N

84.1 13.9 0.4 0.2 0.4 0.8 0.1 0.2

Matrix r 0.7 0.6 0.1 0 0 0.1 0 0.1

84.3 13.8 0.4 0.1 0.3 0.7 0.1 0.2

0.6 0.6 0.1 0 0 0.1 0 0.1

Bulk r 83.4 13.8 0.7 0.2 0.5 0.8 0.1 0.2

800 °C 2 h

Matrix r 0.2 0.9 0.1 0.1 0.2 0.2 0.0 0.0

84.2 13.5 0.7 0.2 0.4 0.6 0.1 0.1

0.5 1.0 0.1 0.1 0.2 0.2 0.0 0.0

Bulk r 84.8 13.3 0.6 0.1 0.3 0.7 0 0.1

however, the core of the particles no longer shows a Cr2O3 type stoichiometry, but instead the core is enriched with Ti and Y, to give a composition close to that of the matrix nanoparticles described above. These results suggest that the initial oxide to form is Cr2O3, and that a Ti-rich shell forms at the interface between the chromium oxide and the matrix on annealing. During consolidation, the core of the particle then transforms to a Y–Ti-based oxide. It is important to note that Fig. 1c–f shows particles >5 nm in diameter; these are also observed occasionally within grains in recrystallized regions. It seems most likely that the >5 nm particles observed within grains initially

Matrix r 0.7 0.5 0 0 0.1 0.1 0 0

1150 °C 100 h

Consolidated

86.2 12.6 0.6 0.1 0.1 0.3 0 0.1

0.7 0.6 0 0 0 0 0 0

Bulk r 85.3 13 0.4 0.1 0.2 0.3 0.1 0.4

Matrix r 0.3 0.2 0 0 0 0 0 0

86 12.7 0.4 0.1 0.1 0.2 0.1 0.4

0.2 0.2 0 0 0 0 0 0.1

Bulk r 87.8 10.9 0.5 0.1 0.1 0.3 0.0 0.1

Matrix r 0.8 1.0 0.1 0.0 0.0 0.1 0.0 0.0

88.4 10.7 0.5 0.0 0.1 0.2 0.0 0.0

0.7 1.0 0.1 0.0 0.0 0.1 0.0 0.0

formed at grain boundaries, and that the grain boundaries moved during recrystallization. 3.3.2. EPFL-E-Fe2Y In EPFL-E-Fe2Y, while the majority of grain boundaries exhibited no solute segregation (as shown in Fig. 9a), several instances of solute segregation to grain boundaries in the MA powder were observed. An example of such a boundary is shown in Fig. 9b. In this instance, an extended section of the boundary is enriched in both Cr and O. the segregation extends over a width of 5 nm, with the Gibbsian excess values for Cr being much higher than

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Table 3 APT measurements of the “matrix-corrected” composition, number density and volume fraction of nanoparticles in the EPFL-E-Y2O3 alloy at all stages of processing (1–5). Element

After MA

800 °C 10 mins

Degassed 800 °C 2 h

Consolidated 1150 °C 4 h 2 MPa

Annealed 1150 °C 100 h

Fe Cr Ti Y O W C N

– 45.22 ± 7.38 16.03 ± 2.18 4.46 ± 0.69 31.82 ± 4.41 0.94 ± 0.19 0.56 ± 0.03 0.81 ± 0.08

– 43.22 ± 3.13 20.46 ± 0.85 4.10 ± 0.38 28.62 ± 2.28 1.35 ± 0.07 0.46 ± 0.14 1.48 ± 0.06

– 51.77 ± 3.37 15.73 ± 1.70 4.82 ± 0.22 23.37 ± 2.08 1.75 ± 0.26 0.41 ± 0.01 1.80 ± 0.11

– 47.80 ± 0.21 12.75 ± 0.55 7.30 ± 0.85 25.55 ± 0.79 0.49 ± 0.04 0.52 ± 0.19 5.35 ± 0.98

– 52.44 ± 0.64 11.92 ± 0.31 8.04 ± 0.2 23.96 ± 0.35 0.23 ± 0.11 0.59 ± 0.03 2.09 ± 0.11

No. density (x1023 m2)

27.3 ± 1.3

21.3 ± 3.2

5.7 ± 0.5

4.4 ± 0.1

3.8 ± 0.4

Volume fraction (%)

0.37 ± 0.02

0.39 ± 0.08

0.21 ± 0.05

0.13 ± 0

0.11 ± 0.04

oxides in the selected analysis volumes). Larger oxide particles are also present at the grain boundaries in the consolidated material. The composition of these >5 nm oxide particles have a similar ratio of Y:Ti:Cr to those in EPFL-E-Y2O3, but in this case the particles contain higher levels of N (at the expense of O). A proxigram of a typical particle outlining the composition is shown in Fig. 10. 4. Discussion 4.1. Validity of the APT technique for the analysis of oxide nanoparticles

Fig. 4. (a) Size distribution of nanoparticles isolated using the maximum separation method in the EPFL-E-Y2O3 material at all stages of processing (1–5). (b) Graphical representation of the matrix corrected nanoparticle compositions at each stage of processing.

that of O, as shown in Table 7. As was the case in EPFL-EY2O3, no large oxide particles were observed in the powder immediately after MA. The Gibbsian excess values for a typical grain boundary in the consolidated material are also included in Table 7. In this case, the grain boundary segregation behavior is very similar to the segregation observed in EPFL-E-Y2O3, with enrichment of Cr, W and C (the higher levels of Y, Ti and O are due to the presence of unavoidable grain boundary

Before discussing the evolution of the nanoparticles and grain boundary oxides in detail, the accuracy of the composition measurements obtained using APT is reviewed. Section 2 compared composition measurements of Y–Ti-based oxide nanoparticles using TEM-based techniques and APT from published literature. There is some disagreement about the composition of the particles measured by various techniques, and it has been previously noted that there are discrepancies in composition measurements made by APT [26]. This may be due to differences in the materials analyzed in each study, but it now appears more likely that the differences in measured composition are a result of differences in the APT data analysis. Proximity histograms [42] have been used to estimate nanoparticle composition [43], but it is more common to use cluster finding algorithms based on the maximum separation method [35] for example [10,23–26,44]. Miller et al. often use the “envelope” method [34] to extract compositional information from the nanoclusters [23–25] whereas this study uses the double maximum separation method. The measured composition is strongly dependent on the parameters required by the software [37] and these differences must be considered when comparing results. Local magnification effects also must be taken into account, as they can affect the measured composition of the nanoparticles. Field ion microscopy experiments have shown that the nanoparticles in ODS steels evaporate preferentially to the matrix [27], and this causes ion trajectory

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Fig. 5. Atom maps showing the distribution of Cr, Y and Ti–O ions in EPFL-E-Fe2Y: (a) stage 1, powder after MA; (b) stage 4, consolidated material.

Table 4 APT measurements of bulk and matrix compositions for processing stages 1 and 4 in EPFL-E-Fe2Y. Bulk measurements are averaged from all atoms in at least three analysis volumes containing 30 M ions+. Matrix composition measurements are estimated by removing the particles using the maximum separation method and only including areas away from grain boundaries exhibiting solute segregation. Element

Fe W C Si Cr Y O Ti N

Fe2Y powder

Consolidated Fe2Y

Bulk

Error

Matrix

Error

Bulk

Error

Matrix

Error

67.66 1.99 0.10 0.03 16.81 0.99 0.79 0.37 1.37

1.6 1.5 0.0 0.0 0.3 0.5 0.2 0.6 0.6

77.84 0.58 0.17 0.04 15.85 0.72 1.64 0.39 1.06

0.2 0.01 0.01 0.27 0.41 0.17 0.49 0.53

83.97 0.47 0.06 0.25 14.10 0.11 0.23 0.21 0.50

1.6 1.5 0.0 0.0 0.6 0.0 0.2 0.6 0.3

84.59 0.48 0.07 0.24 13.85 0.05 0.08 0.11 0.43

0.07 0.01 0.01 0.02 0.08 0.00 0.00 0.01 0.03

aberrations during APT data acquisition. This results in an increase in atomic density in the vicinity of the particles [38]. This is contrary to the behavior seen by of Kluthe et al. for the MgO–Ag system, where lower atomic densities

are found in the vicinity of MgO particles in an Ag matrix [45]. Kluthe et al. suggest that for the MgO–Ag material system, the difference in evaporation field between the particle and the matrix (the particles in this case require a

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Table 5 APT measurements of the “matrix-corrected” composition, number density and volume fraction of nanoparticles in EPFL-E-Fe2Y at stages 1 and 4 of processing. Element

After MA

Consolidated 1150 °C 4 h 2 MPa

Fe Cr Ti Y O W C N

– 39.77 ± 1.09 13.46 ± 2.03 3.51 ± 0.88 34.03 ± 5.37 1.85 ± 1.05 0.78 ± 0.28 3.21 ± 1.02

– 49.42 ± 0.5 12.31 ± 0.17 7.22 ± 0.13 23.95 ± 0.24 0.45 ± 0.07 0.51 ± 0.04 5.97 ± 0.18

No. density (x1023 m2)

77.8 ± 1.8

21.3 ± 3.2

Volume fraction (%)

0.26 ± 0.05

0.13 ± 0.03

higher evaporation field than the matrix) could account for the substoichiometric oxide composition obtained in that study, by the “tearing off” of part of the oxide particle or because of the difficulty in forming positive oxygen ions [45]. An in-depth study on the effect of the trajectory aberrations on the composition of the nanoparticles has been carried out on Y–O-based nanoparticles in an ODS steel [46]. It was shown that when the evaporation field is systematically decreased (by increasing the laser power) the M:O ratio tends towards 1:1, and the Cr content of the clusters decreases [46]. While there may still be difficulty in forming positive oxygen ions in the ODS steel, the trajectory aberrations are not believed to affect the composition of the nanoparticles significantly. This is partly due to the “matrix correction” applied to the composition measurements. The high levels of Cr associated with the nanoparticles that are found in similar ODS steels have been attributed to a core– shell structure [26,27]. While a Cr shell cannot be resolved for the smallest particles in this study, the existence of such a shell is the most likely explanation for the observed levels of Cr, given that Cr is almost insoluble in some Y–Ti oxides [47].

Fig. 7. Grain boundary analysis of EPFL-E-Y2O3 2-D density profile taken from a 2 nm thick slice through an APT analysis volume (grain boundaries are indicated by arrows): (a) stage 1, powder after MA; (b) stage 4, consolidated material. 1-D concentration profile acquired from the cylinder indicated showing the distribution of Cr, W and C across the boundary: (c) stage 1, powder after MA; (d) stage 4, consolidated material.

Similar trends have also been observed for Y–Ti–O nanoparticles [36]. Table 8 is reproduced from Ref. [36], and shows the matrix corrected composition of oxide nanoparticles in a very similar material to that investigated in this study. The EPFL-PA-Y2O3 material has nominally the same composition as the EPFL-E-Y2O3 material, but was manufactured using a pre-alloyed powder rather than elemental powders. Table 8 shows that the composition of the nanoparticles after APT analysis using a relatively low laser power (0.3 nJ) shows the same Y + Ti:O ratio as the

Fig. 6. (a) Size distribution of nanoparticles isolated using the maximum separation method in EPFL-E-Fe2Y at stage 1 and (b) stage 4 of processing. (c) Graphical representation of the matrix corrected nanocluster compositions at stages 1 and 4 of processing.

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Table 6 Gibbsian interfacial excess values and “matrix corrected” grain boundary oxide particle compositions from randomly orientated grain boundaries in EPFL-E-Y2O3 at each stage of processing (1–5). C values shown in atoms nm2 and particle compositions are given in at.%. Element

800 °C 10 mins

After MA 2

C (at/nm ) Particle comp. Cr W C Y Ti O

0 0 0 0 0 0

2

C (at/nm ) Particle comp.

None 17.5 ± 2 observed 4.9 ± 0.5 1.1 ± 0.3 0 0.8 ± 0.2 3.6 ± 0.6

37.9 ± 2.0 1.0 ± 0.1 0.3 ± 0.1 0.4 ± 0.1 7.4 ± 2.3 51.1 ± 1.0

800 °C 2 h

1150 °C 100 h

Cons.

2

C (at/nm ) Particle comp.

C (at/nm ) Particle comp.

C (at/nm2) Particle comp.

4.5 ± 2.5 0.6 ± 0.2 0.1 ± 0.1 0.0 ± 0.1 0.1 ± 0.1 0.1 ± 0.1

2.9 ± 0.9 0.5 ± 0.2 1.0 ± 0.7 1.9 ± 1.4 1.4 ± 1.4 2.1 ± 0.9

3.7 ± 1.0 1.2 ± 0.1 0.3 ± 0.03 0.0 ± 0.1 0.1 ± 0.0 0.1 ± 0.0

material analyzed using a much higher laser energy of 1 nJ. The Cr level, however, is decreased, and this is indicative of Cr forming a shell around the nanoparticles. This is because when the laser energy is increased, the effective evaporation field of the matrix is decreased. This reduces the difference between the actual evaporation field during analysis and the lower evaporation field required for the evaporation of the nanoparticles, and thus lessens the lensing effect on the ion trajectories. It is only since the recent increase in the use of laserpulsed APT that bulk insulating materials have been analyzed successfully using this technique, e.g. Refs. [48,49]. The expected stoichiometry is not always achieved [45,49], and the authors cite many possible reasons for this. It is possible that effects such as off-pulse ion evaporation, poor thermal conductivity resulting in overlapping peak tails or multiple ion evaporation events could lead to oxygen substoichiometry in the nanoparticles investigated in this study but it is unlikely that any of these effects are able to explain such an apparent loss of oxygen. Firstly, Ref. [46] shows that bulk Y2O3 can be analyzed giving the expected stoichiometry and furthermore, Ref. [10] shows that Y–O nanoparticles with a stoichiometric Y:O ratio of 2:3 are observable by APT. Analysis of bulk Y2Ti2O7 and Y2TiO5 in Ref. [36] also do not show any apparent loss of oxygen. In this study, grain boundary oxides with near stoichiometric Cr2O3 composition are also observed (see Fig. 8). The combination of these results indicates that there is no reason to assume that in this study there has been a systematic oxygen loss due to the APT analysis. The interpretation of the data presented in Section 4.2 (discussed below) also suggests that a substoichiometric oxide composition is viable. 4.2. Evolution of nanoparticles in EPFL-E-Y2O3 Regardless of whether the starting powder contains Fe2Y or Y2O3 the results presented in Tables 3 and 5 show that the nanoparticles do not appear to undergo any major compositional change during processing. The evolution of the nanoparticles will now be discussed primarily in relation to EPFL-E-Y2O3. In this material, visual inspection of the changes in composition shown in Fig. 4b highlights the similarities in the composition of the nanoparticles at

23.2 ± 2.2 1.8 ± 0.2 0.1 ± 0.1 0.6 ± 0.3 22.9 ± 3.3 49.4 ± 1.6

2

27.1 ± 6.2 1.5 ± 0.1 1.6 ± 0.2 7.0 ± 2.6 20.2 ± 0.1 42.4 ± 6.2

11.2 ± 0.5 2.8 ± 1.5 1.0 ± 0.5 11.8 ± 1.0 23.9 ± 3.9 44.9 ± 0.6

each stage, but Table 3 shows that there is a slight decrease in oxygen following the annealing of the powder, and an increase in Y content after consolidation. The O level in the nanoparticles directly after MA is initially slightly higher than in subsequent stages, but this must be considered in the wider context of the evolution of the whole microstructure. Not only is the oxygen level in the nanoparticles higher, but the level in the matrix (shown in Table 2) is also higher. As the grain structure develops, so does the formation of grain boundary oxides. In Table 3 we observe the volume fraction occupied by the nanoparticles and the number density of nanoparticles to decrease as the material is processed. This is an indication that solutes associate preferentially with larger grain boundary oxides. Fig. 9 suggests that formation of the grain boundary oxides is preceded by segregation of oxygen to grain boundaries. Because of the high density of nanoparticles after MA, this suggests that the lower oxygen concentration in the nanoparticles at the latter stages of processing is more stable. We conclude from these results that the substoichiometric quantity of O in the nanoparticles is not a result of there being insufficient O available from the matrix at the nucleation stage. We must assume therefore that the amount of O in the nanoparticles is constrained by effects relating to the matrix in which the nanoparticles form. On a basic level, the precipitation of the nanoparticles is driven by the low equilibrium solubility of O in the matrix. First and foremost it is an oxidation reaction that causes a nanoparticle embryo to form, driven by the change in free energy associated with that reaction. Table 9 shows the formation enthalpies for many of the possible stable oxides. The formation of complexes with Y and Ti is expected given that these elements have a high affinity for oxygen, and amongst these known stable oxides, the formation of Y2Ti2O7 is considered most likely. However neither the “non-stoichiometric” M:O ratio nor Y:Ti ratio in the nanoparticles is consistent with this, nor is it consistent with the rock-salt structure of the nanoparticles proposed by Hirata et al. [21]. The kinetics of the nanoparticle nucleation will depend on the concentrations and diffusivities of the oxide-forming species, and the energy barrier for nucleation. The nucleation barrier can be considered as a combination of a strain

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Fig. 8. Proximity histograms showing the distribution of solute elements from a 1% Ti isoconcentration surface representing the grain boundary oxide particle interface for a representative particle in EPFL-E-Y2O3 after (a) stage 2, (b) stage 3 and (c) stage 4 of processing.

energy term corresponding to the change in volume free energy associated with the nanoparticle, and also the extra energy associated with the formation of a metal–oxide interface. In highly deformed alloys, second phase particles are assumed to nucleate heterogeneously on dislocations [50], and there will be an abundance of such nucleation sites introduced by the MA. Hirata et al. have shown that a defective rock-salt structure can be fully coherent with the matrix, and suggested that this gives rise to a very low interface energy [21]. This is likely to encourage the formation of non-stoichiometric nanoparticles over the more thermodynamically stable Y2Ti2O7. Ribis and de

Carlan have shown that after annealing at 1300 °C the nanoparticles have a structure corresponding to Y2Ti2O7, and identified a well-defined orientation relationship that means that the nanoparticle–matrix interface is also coherent with the matrix [20]. We therefore propose that the formation of Y2Ti2O7 nanoparticles requires higher temperatures, because as the temperature is increased, the elastic strain energy term is reduced by an increase in solute diffusion. The change in structure is more restricted at the lower temperatures used in this study. Fig. 3 and Table 3 show that the nanoparticles are formed at relatively low temperatures during MA, and evolve with a near-same

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Fig. 9. Grain boundary analysis of two different boundaries in EPFL-E-Fe2Y powder after MA (stage 1). (a and c) 2-D density profile taken from a 2 nm thick slice through an APT analysis volume. Grain boundaries are indicated by arrows. (b and d) 1-D concentration profile acquired from the cylinder indicated showing the distribution of Cr, W and C across the boundary.

Table 7 Gibbsian interfacial excess values and “matrix corrected” grain boundary oxide particle compositions from the grain boundaries indicated in Fig. 9 in EPFL-E-Fe2Y powder after MA (stage 1) together with values from EPFL-E-Fe2Y consolidated material (stage 4). C values shown in atoms nm2 and particle compositions are given in at.%. Element

Fe W C Si Cr Y O Ti N

After MA

Consolidated

C (atoms/ nm2)

Particle comp.

C (atoms/ nm2)

GB particle composition (at.%)

0 0 0 0 0 0 0 0 0

None observed

0.0 ± 0.0 1.7 ± 0.4 2.8 ± 1.2 2.4 ± 0.6 18.7 ± 1.3 3.3 ± 0.9 12.1 ± 2.7 6.6 ± 1.9 1.6 ± 0.1

0.0 ± 0.0 0.3 ± 0.1 0.1 ± 0.1 0.6 ± 0.3 33.7 ± 2.1 1.4 ± 0.6 21.7 ± 1.0 25.4 ± 1.7 16.6 ± 0.7

composition as the particles in the consolidated material after annealing at 800 °C. We conclude that, at these temperatures, a transformation to Y2Ti2O7 is improbable. Combining the findings of Hirata et al. and the findings in this study gives a greater understanding of the formation mechanism of the nanoparticles. It appears that the nanoparticles form with the basis of a TiO type structure. This is consistent with the surface oxidation behavior of Ti in low oxygen environments, where the reaction proceeds with the

Fig. 10. Proximity histogram showing the distribution of solute elements from a 1% Ti isoconcentration surface representing the grain boundary oxide particle interface for a representative particle in EPFL-Fe2Y after stage 4 of processing.

formation of TiO in preference to the more common TiO2 [51]. Hirata et al. noted that the rock-salt-type structure is very accommodating to vacancies, and can adopt a wide range of chemistries [21]. This allows an explanation for the non-stoichiometric Y:Ti ratio, which is likely to relate

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Table 8 “Matrix-corrected” cluster compositions for nanoparticles in the EPFL-EY2O3 material after analysis APT analysis using a low laser energy of 0.3 nJ and a high laser energy (1 nJ). Error calculated as standard error from all clusters in an analysis volume of 10 M ions. Element

0.3 nJ laser (at.%)

1 nJ laser (at.%)

Cr Ti Y O

44.3 ± 0.9 21.1 ± 0.4 5.6 ± 0.2 24.3 ± 0.4

30.3 ± 0.5 23.8 ± 0.2 8.9 ± 0.1 28.8 ± 0.2

Y + Ti:O

1.1

1.1

to how predisposed these cations are to being found in the 2 + valence state. Y2+ cations are uncommon, but have been observed experimentally by Johnson et al. [52]. Kang and Bernsein also see abundant evidence of Y2+ in the formation of yttrium oxide particles using pulsed laser vaporization; however, this is only in the case when the oxygen content of the seed gas used in the vaporization process is close to zero [53]. For charge balance to be maintained in the nanoparticles, both Y and Ti would need to be found predominantly in the 2 + valence state. As Ti2+ ions are more readily formed than Y2+ ions, a dominance of these ions in the nanoparticles is unsurprising, and we propose that Y ions take the place of Ti with the assistance of the high concentration of vacancies in the system. Table 2 shows that there is no significant change in the matrix composition between the degassing stage and the consolidated material (or the consolidated material annealed at 1150 °C for 100 h). This indicates that growth of the nanoparticles is complete by the end of the degassing stage, and any subsequent size change in the nanoparticle distribution occurs via a coarsening mechanism. A previous study on the same material has shown that at 1200 °C the nanoparticles coarsen according to classical coarsening kinetics [10], and the broadening of the particle size distribution after annealing at 1150 °C for 100 h is consistent with this. However, the changes in the volume fraction of the nanoparticles shown in Table 3 do not follow the trends predicted by classical nucleation, growth and coarsening theory. If this were the case, the volume fraction

of nanoparticles would steadily increase during the growth stages, and then plateau during the coarsening stages. Instead, there is an overall decrease in volume fraction of nanoparticles. This is assumed to be due to the formation of the grain boundary oxides occurring after the nanoparticles, and reducing the amount of solute available to the nanoparticles. This reasoning is confirmed by the EPMA analysis shown in Table 2, where the results suggest that a sizable proportion of Cr, Y, Ti and O are outside the typical volume of analysis for APT analysis. The implication of these findings is that the nanoparticles are less stable than the larger grain boundary oxides. The formation of these larger oxides will be addressed below. 4.3. Effect of mechanical alloying with Fe2Y vs. Y2O3 on nanoparticles The results presented above show that the use of Fe2Y instead of the more conventional Y2O3 to incorporate Y into the alloy has no significant effect on the oxide nanoparticle distribution. The composition, size distribution and number density of Y–Ti–O-based nanoclusters are virtually identical in each alloy, both for the powder directly after MA and in the consolidated material. This leads to two conclusions. Firstly, the majority of the oxygen required to generate the oxide nanoparticle dispersion must be incorporated from the atmosphere and/or surface oxide during processing, and does not need to be deliberately included in the initial starting powder. Secondly, the way in which Y and O is incorporated into the powder during MA is not a controlling factor in the generation of the oxide nanoparticle dispersion. This is contrary to the findings of Oksiuta et al., who suggested that less dense population of Y–O nanoparticles with 5 nm radius are formed when mechanically alloying with Fe2Y [32]. One possible reason for this difference is the added milling time (50 h in this study compared to the 20 h used by Oksiuta et al.). Increasing milling time allows more oxygen to be incorporated, thus increasing the overall amount of O available in the matrix, and may promote the formation of the Y–Ti–O type nanoclusters.

Table 9 Enthalpy of formation of oxides (data for 298 K) from Ref. [58] unless otherwise stated.

4.4. Formation and evolution of grain structure and grain boundary oxides

Stoichiometric oxide phase

M:O ratio

Enthalpy of formation DHf per mole of oxide

Enthalpy of formation DHf per mole O

(Least ve) CrO2 Cr2O3 TiO2 Ti3O5 YCrO3 Ti2O3 TiO Y2Ti2O7 (Most ve) Y2O3

1:2

583

292

2:3 1:2 1:1.67 2:3 2:3 1:1 1:1.75 2:3

1130 944 2457 1493[50] 1522 543 3874[51] 1907

377 472 491 497 507 543 553 636

The existence of larger grain boundary oxides could have implications for the mechanical properties of the alloys, and their formation has been shown to affect the evolution of the oxide nanoparticles. To understand the mechanism of formation of these oxides, both the evolution of the grain structure and the chemistry of the grain boundaries during processing needs to be understood. While there is no significant difference in the oxide particle dispersions in consolidated EPFL-E-Fe2Y compared to EPFL-E-Y2O3, there is a slight increase in grain size. It is possible that this subtle difference is related to the strength of the initial ceramic Y2O3 particles compared to the Fe2Y

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intermetallics. Comparing TEM images of the microstructure for consolidated EPFL-E-Y2O3 in Fig. 1c and consolidated EPFL-E-Fe2Y in Fig. 2b there is a higher density of dislocations in EPFL-E-Y2O3. These dislocations could be generated by continuous impacts with the higher strength particles during MA. A lower initial dislocation density in the MA powder after MA with Fe2Y is likely to speed up the recovery and recrystallization process, and allow the grains to grow slightly larger before the formation of the grain boundary oxides. This means that effectively the grain development seen in EPFL-E-Fe2Y MA powder has progressed further than in EPFL-E-Y2O3 at the same stage of processing. In the discussion that follows, it is assumed that the grain development in both alloys proceeds in a similar way, and that EPFL-E-Fe2Y at stage 1 of processing is “ahead“ of EPFL-E-Y2O3 at the same stage. However, further study on the relative dislocation densities across a number of different batches of material is required to establish whether the differences in the grain structure observed here is truly related to the initial starting powders, or whether it is just inherent variation between one batch of material and the next. Fig. 1 shows that the grain structure in EPFL-E-Y2O3 changes dramatically during the initial stages of processing. However, once formed, the submicron grain structure of the consolidated material is very stable, with no significant change after annealing for 100 h at 1150 °C. This is a common feature of ODS materials. Bhadeshia attributes the high temperature stability to grain boundary pinning by the triple-junctions of the grain boundaries themselves [54]. The high density of oxide particles almost certainly also contributes to stabilizing the grain boundaries. The interfacial excess of W, Cr and C to grain boundaries shown in Tables 6 and 7 demonstrates that in both the annealed powder and consolidated material, the grain boundaries are very accommodating to solute elements. However, Figs. 7a and 9a show that in the powder immediately after MA, no segregation of solutes corresponding to grain boundaries is observed. The oxide particles are known to suppress recrystallization below 800 °C [55]. It is therefore assumed that recrystallization has not occurred at this stage. The “grains” observed in Figs. 1b and 2a are more likely to be the result of polygonization and hence are mainly subgrains separated by low-angle boundaries comprising an ordered array of dislocations. In the EPFL-E-Fe2Y powder directly after MA, pockets of O and Cr enrichment are occasionally observed at specific boundaries (assumed to be high-angled boundaries), as shown in Fig. 9b. If excess O from the matrix migrates to high-angle boundaries, this would help to explain the much higher levels of O in the bulk directly after MA that are indicated in Table 2. Initially, the crystal structure in the material is heavily deformed, therefore a higher proportion of O is still present in the matrix than after annealing. The high Cr and O levels seen at the boundary in Fig. 9b are likely to represent the initial stages of formation of grain boundary oxide particles. The evolution of the

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core–shell structure and composition of these grain boundary oxides in EPFL-E-Y2O3 can be seen in Fig. 8. Despite the enthalpy of formation of Y and/or Ti oxides being more favorable (as shown in Table 9), the initial oxide to form at the grain boundaries is a Cr-based oxide. As the O level is so far above the equilibrium saturation limit of a few ppm [56], the reaction pathway to precipitate out the excess O will be controlled by the kinetics of oxidation. Due to the high matrix level of Cr, the fastest oxide to form will be Cr2O3 because with 14% Cr, the distance over which Cr must diffuse becomes negligible. This is analogous to the results of Hin and Wirth, who investigated precipitation in Fe–Y–O and Fe–Ti–O alloys using kinetic Monte Carlo methods. There, the authors found that prior to the more stable Y or Ti oxides forming,small Fe2O3 precipitates formed [57]. As Cr has a greater affinity for O than Fe, the formation of Cr2O3 by a similar mechanism is expected in the alloys studied here. Table 10 shows the diffusion coefficients of each of the solute elements in body-centered cubic (bcc) Fe at the degassing temperature. Based on diffusion coefficients alone, a Ti-based oxide would be expected, and at the second stage of characterization (i.e. after 800 °C for 10 min) the Ti-shell is already observed. If a uniform concentration gradient between the nanoparticle and the matrix is assumed, the diffusion distance of Y is 6 nm after annealing for 10 min whereas O can diffuse >20 lm in the same time. At 800 °C the diffusion of Y is orders of magnitude slower than the other clustering elements, therefore it is logical that the higher temperature and pressure during consolidation is required to allow Y to diffuse to these particles. This mechanism of formation means that after the formation of the grain structure, avoiding the formation of the grain boundary oxides is likely to be possible only by minimizing the O uptake in the powder. It was shown that this is difficult to control, as MA with Fe2Y did not affect the oxygen content of the MA powder significantly. In summary, the results presented above indicate that the formation and evolution of the grain boundaries and associated oxide particles proceeds in the following way: 1. In MA powder initially the material lacks any form of well-defined grain structure, and the boundaries observed between nanometer-scaled crystalline regions do not exhibit solute segregation.

Table 10 Diffusion coefficients for particle forming species in bcc-Fe at 800 °C [58] (Ti) and [59]. Element

Diffusion coefficient (m2 s1)

O Ti Cr Y

5.7  1013 2.5  1015 7.5  1016 2.8  1020

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2. Oxygen migrates to the grain boundaries to form pockets of O- and Cr-rich regions at the boundaries (observed occasionally in the EPFL-E-Fe2Y MA powder). 3. The oxygen-rich regions attract Ti and later Y to form second phase particles at the interface. 4. Simultaneously, as the material is heated during degassing and subsequent stages, dislocations order to form grain boundaries between crystalline regions, and segregation of W and Cr is favorable as this reduces misfit strain.

5. Conclusions The effects of individual stages of ODS alloy processing has been investigated by systematic characterization of the material by APT and TEM. The effect of MA with Fe2Y compared with Y2O3 to incorporate Y into the system during mechanical alloying has also been investigated. This has provided some key insights into the formation of nanoparticles and larger oxides in ODS steels.  The formation of Y–Ti–Cr–O nanoparticles in a ferritic ODS alloy occurs during the MA alloy process, and the initial particles maintain a similar ratio of Cr:Ti:Y:O throughout the processing.  As the same nanoparticle dispersion is observed after MA with Fe2Y and Y2O3 it is clear that the majority of oxygen in the nanoparticles is incorporated from the atmosphere during MA.  There is a slight reduction in O content of the nanoparticles after the degassing stage.  After consolidation, the concentration of Y in the nanoparticles increases from 4% to 7%.  The grain structure evolves concurrently with the nanoparticles. Initially no segregation to grain boundaries is observed, but C, W, Cr and oxide particles >5 nm are observed after stage 2 of processing.  At the MA stage there is a high proportion of Y, Ti and O in the matrix. These proportions decrease after annealing the powder for 10 min at 800 °C, and are reduced to trace levels by the end of the degassing stage (800 °C, 2 h anneal).  The formation of grain boundary oxides is shown to influence the proportion of Y, Ti and O available to the nanoparticles. The M:O ratio of the nanoparticles is therefore not defined by the O available in the system.  Oxide particles >10 nm observed at grain boundaries have a Ti–Cr–O-rich shell, and contain no Y before consolidation, but have similar composition to the nanoparticles at the core after consolidation.  An intermediate Cr2O3 grain boundary oxide is observed in the powder after annealing at 800 °C for 10 min and 2 h. From this it can be concluded that the size and composition of the nanoparticles cannot be controlled by thermal treatments before or after consolidation at industrially

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