Influence of ball-milling and annealing conditions on nanocluster characteristics in oxide dispersion strengthened steels

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Acta Materialia 60 (2012) 7150–7159 www.elsevier.com/locate/actamat

Influence of ball-milling and annealing conditions on nanocluster characteristics in oxide dispersion strengthened steels M. Laurent-Brocq a,⇑, F. Legendre a, M.-H. Mathon b, A. Mascaro a,1, S. Poissonnet a, B. Radiguet c, P. Pareige c, M. Loyer a, O. Leseigneur a a

CEA, DEN, Service de Recherches de Me´tallurgie Physique, F-91 191 Gif-sur-Yvette, France b CEA-CNRS, Laboratoire Le´on Brillouin, 91 191 Gif-sur-Yvette, France c Groupe de Physique des Mate´riaux, UMR CNRS 6634, Universite´ et INSA de Rouen, 76801 Saint Etienne du Rouvray, France Received 11 July 2012; received in revised form 6 September 2012; accepted 10 September 2012 Available online 13 October 2012

Abstract The characteristics of strengthening nanoclusters (NCs) have a major influence on the mechanical properties of oxide dispersion strengthened (ODS) steels. To determine how to control NC formation, ODS powders are synthesized in different ball-milling and annealing conditions, then characterized by electron probe micro-analysis and small angle neutron scattering. During ball-milling, reactants are dissolved into the metallic matrix until a Ti, Y and O solid solution is formed and then NC nucleation begins. Nucleation is greatly enhanced during the first minutes of annealing at 800 °C without any coarsening afterwards. The intensity and temperature of ball-milling influence this mechanism and thus the characteristics of the formed NC, whereas the nature of reactants, for a given composition, has no impact on NC size and volume fraction. Consequently, to promote the formation of fine and dense dispersion of NC, two main modifications to the usual process are proposed: (i) perform a long and/or intense ball-milling with a limited overheating, and (ii) anneal the as-milled powder at 800 °C before performing the thermo-mechanical treatment. Ó 2012 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Oxide dispersion strengthened steels; Nanoparticles; Mechanical milling; Diffuse neutron scattering

1. Introduction The temperatures and the neutron doses in the core of the generation IV and fusion nuclear reactors will be higher than those in generation II and III reactors. To realize engineering solutions for such severe conditions, new structural materials are needed. Oxide dispersion strengthened (ODS) steels are promising candidates. They consist of a ferritic or martensitic (F/M) matrix, which exhibits a low swelling under irradiation, containing dispersed nano-oxides, which ⇑ Corresponding author. Present address: University of Namur, FUNDP, Research Centre in Physics of Matter and Radiation (PMR),

E-mail address: [email protected] (M. Laurent-Brocq). Present address: CEA, DEN, DMN, SRMA, Laboratoire d’Analyse Microstructurale des Mate´riaux, 91 191 Gif-sur-Yvette, France. 1

reinforce the poor creep strength of the F/M matrix at high temperatures [1–3]. However, not all ODS steels have equivalent creep strength [1]. This is partly due to the characteristics of the nano-oxides dispersion. Indeed, according to the particle–dislocation interaction theory, the highest deformation threshold stress is obtained for the smallest inter-particle distance, which is equivalent, at a given volume fraction of particles, to the smallest particle radius [4]. By characterizing the MA 957 ODS steel, Sakasegawa et al. [5] have identified two categories of nano-oxides whose nature and composition are size-dependent: (i) from 1 nm to 15 nm, Y-, Ti- and O-enriched nanoclusters which have been extensively characterized by atom probe tomography [6–11]; and (ii) from 15 nm to 35 nm, stoichiometric Y2Ti2O7 precipitates mainly identified by transmission electron microscopy (TEM) [12–14]. Due to their smaller

1359-6454/$36.00 Ó 2012 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.actamat.2012.09.024

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size, nanoclusters are the most efficient type of reinforcement for tensile and creep properties. This statement has been confirmed experimentally by Hoelzer et al. [15] through the comparison of two alloys: the first contains numerous Y2Ti2O7 oxides of 20 nm and few nanoclusters of 5 nm heterogeneously dispersed, while the second is reinforced by a dense and homogeneous dispersion of 5 nm nanoclusters and by a dense but heterogeneous dispersion of 1 nm nanoclusters. Indeed, the second alloy does exhibit greater tensile properties. The oxide dispersion also improves the swelling resistance and it is more efficient if the dispersion is fine and dense. Indeed Hsiung et al. have observed by high-resolution TEM in an irradiated ODS steel that helium bubbles are formed preferentially at the metal/oxide interface, and that the smaller the nano-oxides are, the smaller the helium bubbles are [16]. So, to obtain creep strength and maximum swelling resistance in ODS steel, the challenge is to find a process through which the formation of a dense dispersion of nanoclusters is promoted and the coarsening of those nanoclusters into stoichiometric oxides is prevented. The standard process to creating an ODS material consists of ball-milling Y2O3 and pre-alloyed powders and then performing thermomechanical treatments [4,17]. The commonly proposed formation mechanism is that Y2O3 is dissolved into the metallic matrix during ball-milling and then nano-oxides are formed during subsequent thermomechanical treatment. However the evidence for this mechanism is based solely on characterization of the final material, at the end of the multi-step process [9,12,18], while there are very few insights into the very beginning of nano-oxides’ formation [19–21]. Moreover, the process depends on many parameters. For example, it was shown that a very pure ball-milling atmosphere, which limits the powder oxidation, results in a thinner nano-oxides dispersion [22,23]. In Ref. [21], nanoclusters were detected by atom probe tomography in the as-milled powder, possibly because of specific milling conditions. The thermo-mechanical treatment temperature was also proven to impact the final size of nano-oxides [15,24]. Still, there are many parameters which may be varied, especially during the milling stage, whose effects on the final microstructure remain unknown. As a consequence it has not yet been possible to precisely monitor the nano-oxides’ characteristics as a function of production parameters. To address these issues, this work focuses on the very first steps of nano-oxide formation, which are ball-milling and annealing. The objective is to identify the formation mechanisms and to determine how parameters influence the nano-oxides’ characteristics. Toward this goal, an ODS steel of one given composition was synthesized through ball-milling in varying conditions (reactants, duration, intensity and temperature) and annealed. As-milled and as-annealed powders were then characterized by electron probe micro-analysis (EPMA), small angle neutron scattering (SANS) and atom probe tomography (APT).

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2. Experimental 2.1. Synthesis The reactants Fe–14 Cr–2 W, Fe–14 Cr–2 W–0, 25 Ti (wt.%), YFe3 and Fe2Ti were prepared in an electromagnetic induction furnace. It was checked by EPMA that O content is lower than 0.6 wt.% in Fe2Ti ingots, that Ti content is higher than 0.2 wt.% in FeCrWTi and that no oxygen is detected in FeCrW and FeCrWTi ingots. Commercial Y2O3 and Fe2O3 powders with a size of few microns were used as well as a nanometric (10 nm) Y2O3 powder. All the reactants were mixed in different combinations so as to consistently achieve the same nominal composition: Fe–14 Cr–2 W–0.25 Ti–0.2 Y–0.05 O (wt.%), or equivalently Fe–15.1 Cr–0.6 W–0.29 Ti–0.13 Y–0.17 O (at.%). To manage the ball-milling conditions, a Fritsch P0 mill was chosen. It was used with a WC ball and vial, under a secondary vacuum atmosphere. Milling lasted between 1 and 144 h. The intensity, as defined by Chen et al. [25], is given by: Mb  x  A  f I¼ Mp where Mb and Mp are the mass of the ball and powder respectively, A and x are the amplitude and the pulsation of the mill movement respectively, and f is the frequency of impact between the ball and the powder. All those parameters were fixed except the vibration amplitude, allowing the ball-milling intensity to vary from 1000 to 4000 m s2. Since the mill P0 has only one ball and vibrates only vertically, ball/powder/wall frictions and the subsequent heating are limited. In normal conditions, an average temperature of 30 °C is measured via a thermocouple on the outer side of the vial. However, local overheating cannot be excluded. To increase the temperature when needed, the mill was surrounded by a resistor. A temperature of 150 °C was reached inside the vial. 20 ODS powders were synthesized with different reactants, duration, intensity and temperature of ball-milling; they are listed in Table 1. Ref-B is the reference ball-milled sample; for this sample, reactants are FeCrW + Fe2Ti + YFe3 + Fe2O3, ball-milling duration is 144 h, intensity is 2000 m s2 and temperature is 30 °C. For other samples, one or two parameters differ from this reference state. Annealing was performed under an argon atmosphere at 800 °C once the sample had reached the desired temperature, which takes about 10 min. Then the sample was quickly cooled (taking 3 min to reach 400 °C, and a total of 15 min to reach 100 °C). A or A2 are added to the asmilled sample names to denote that annealing was performed at 800 °C for 5 min or 60 min respectively. 2.2. Characterization techniques Electron probe microanalysis (EPMA) was done with a Cameca SX50, at a 15 kV voltage and a 40 nA current. X-

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Table 1 Milling conditions of the ODS samples. d, I and T are respectively the duration, the intensity and the temperature of the milling. All samples have the nominal composition Fe–14 Cr–2 W–0.25 Ti–0.2 Y–0.05 O (wt.%). Conditions corresponding to the ones of the reference sample are in bold. Name

Milling conditions Reactants

d (h)

I (m s2)

Ref-B B1 B5 B10 B24 B48 B72 B100

FeCrW + Fe2Ti + YFe3 + Fe2O3

144 1 5 10 24 48 72 100

2000 2000 2000 2000 2000 2000 2000 2000

30 30 30 30 30 30 30 30

Re1

FeCrWTi + Y2O3 nm

144

2000

30

Re2

FeCrWTi + Y2O3

144

2000

30

Re3

FeCrWTi + YFe3 + Fe2O3

144

2000

30

I1-B24 I1-B72 I1-B I3-B24 I3-B72 I3-B T-B24 T-B72 T-B

FeCrW + Fe2Ti + YFe3 + Fe2O3

24 72 144 24 72 144 24 72 144

1000 1000 1000 4000 4000 4000 2000 2000 2000

30 30 30 30 30 30 150 150 150

T (°C)

ray maps and weight quantitative measurements were generated. The atom probe tomography (APT) sample preparation was achieved in two steps. First a powder grain was stuck by silver lacquer on a stainless-steel tip with an end radius of several micrometers, using a micromanipulator under an optical microscope. Then the powder grain was machined into a tip-like shape in an SEM-FIB (scanning electron microscope-focused ion beam). 30 keV Ga ions were used to machine the grain. When the tip radius was smaller than 50 nm, the ion energy was decreased down to 2 keV so as to avoid Ga implantation and radiation damage in the part of the sample which would be analysed. Atom probe analyses were performed with Cameca LATAP (laser assisted tomographic atom probe) at 80 K, using fs laser pulses with a wavelength of 515 nm, a pulse repetition rate equal to 2 kHz, and a power setup so as to represent 20% of the standing potential. To analyse the spatial distribution of elements, frequency distribution tests [26] were performed on all solute evaporating species, that is to say on single ions or molecular ions (such as YO or TiO). Small angle neutron scattering (SANS) experiments were performed at Laboratoire Le´on Brillouin, on the PAXY and PAXE instruments. The choice of experimental conditions and the hypotheses for data treatment are detailed and justified by Mathon et al. in Ref. [27]. Two configurations of measurements were used, defined by wavelengths k of 0.6 or 0.9 nm and sample-to-detector distances of 2 or 5 m, covering a total scattering vector (q) range from 0.07 to 1.6 nm1 (q ¼ 4p  sin h, where 2h is k

the scattering angle). This allows for the detection of nanoclusters (NC) of 1 or 2 nm as well as stoichiometric oxides of tens of nanometers. Measurements have been performed at room temperature, under a saturating magnetic field H (=1.7 T) perpendicular to the incident beam, so as to be able to separate the magnetic and nuclear scattering cross-sections. As-milled samples were introduced in a quartz container calibrated to have a thickness of 1 mm parallel to the neutron beam. As-annealed samples were in the form of platelets with a diameter of 8 mm and a thickness from 0.95 to 1 mm. Data were treated with an absolute calibration as defined in Refs. [28,29]. Plexiglas was chosen as the standard sample and for as-milled samples, the contribution of the empty container was subtracted. Fe–14 Cr–2W (wt.%) alloys, which were milled and annealed in the same conditions as ODS powders, were used as reference samples (see Table 2). An average incoherent scattering of 0.1 cm1 was calculated by applying Porod’s law to those matrix samples [30]. It was subtracted from the cross-sections of every sample. At very small angles, cross-sections of all samples follow Porod’s law, which means that the diffusion is due only to interfaces. For higher angles (corresponding to q > 0.4 nm1), the signal is due to precipitates of few nanometers. This part of the signal was quantitatively analyzed. It was assumed that precipitates are isotropic, diluted, spherical, homogeneous, have the same composition and have a Gaussian size distribution. Thus nuclear plus magnetic and magnetic cross-sections can be written as:   dr 1 ð~ qÞ ¼ f  ðDq2mag þ Dq2nucl Þ dX nuclþmag VS R1 hðRÞ  V 2 ðRÞ  F 2 ðR;~ qÞ  dR  0 R1 hðRÞ  V ðRÞ  dR 0 R1   hðRÞ  V 2 ðRÞ  F 2 ðR;~ qÞ  dR dr 1 2 ð~ qÞ ¼ f  Dqmag  0 R 1 dX mag VS hðRÞ  V ðRÞ  dR 0 where h is a Gaussian function defined by an average radius Rm and a width g; f, Dq and h are the volume fraction, the contrast (nuclear or magnetic) and the size distribution of nanometric precipitates respectively; F is the shape factor of a sphere; Vs and V(R) are the volume of the sample and the volume of a sphere of radius R respectively. The fitting parameters are Rm, g and the product f  (Dqmag2 + Dqnucl2) or f  Dqmag2. The average radius and the width of the distribution of nanometric precipitates were deduced from the nuclear plus magnetic cross-sections while the volume fraction was deduced from the magnetic cross-section. The hypothesis was made that nano-oxides are non-magnetic, allowing for a calculation of the magnetic contrast of precipitates with the following definition:   l lm Dqmag ¼ a   V Vm

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Table 2 Characteristics of nanoporosities measured by SANS in as-milled and as-annealed matrix powders. The reactants were FeCrW. d, I, T, R and f are respectively the duration, the intensity, the temperature, the average radius and the volume fraction. Uncertainties are 5% and 15% respectively for R and f. Name

M-B24 M-B72 M-B M-B-A M-B-A2

Milling conditions

Annealing conditions 2

Nanoporosities characteristics

d (h)

I (m s )

T (°C)

T (°C)

d (min)

R (nm)

f (%)

24 72 144 144 144

2000 2000 2000 2000 2000

30 30 30 30 30

– – – 800 800

– – – 5 60

1.8 ± 0.09 1.5 ± 0.08 1 ± 0.05 1 ± 0.05 1.2 ± 0.06

0.3 ± 0.05 0.2 ± 0.03 0.1 ± 0.02 0.1 ± 0.02 0.3 ± 0.05

0

˚ 3; l and lm are the mean where a = 0.2695  1012 cm1 A atomic magnetic moments under a saturating magnetic field of the nanometric precipitates and the matrix respectively; V and Vm are the mean atomic volume of the nanometric precipitates and the matrix respectively. Here, 0 ˚ 3, l = 0 as a consequence of the hypothesis Vm = 11.77 A and lm = (2.2  2.39  CCr)  lB as defined in Ref. [31], where CCr is the chromium content and lB the Bohr magneton. In the following figures only the nuclear plus magnetic cross-sections are plotted. Matrix samples contain a population of nanometric non-magnetic objects. By comparing their nuclear and magnetic cross-sections, those objects were identified as nanoporosities. The as-milled samples have nanoporosity volume fractions of 0.3, 0.2 and 0.1% after 24, 72 and 144 h of ball-milling, respectively. The annealed samples have nanoporosity volume fraction of 0.1 and 0.3% after an annealing at 800 °C of 5 and 60 min respectively (see Table 2). Those nanoporosities volume fractions were subtracted from the non-magnetic precipitates volume fractions of the corresponding ODS samples so as to obtain the volume fraction of nano-oxides.

3. Results 3.1. Ball-milling duration To investigate the formation mechanism of nano-oxides, ball-milling was performed for a number of durations, ranging from 1 to 144 h. The reactants were FeCrW + Fe2Ti + YFe3 + Fe2O3. Ball-milling intensity and temperature were fixed respectively at 2000 m s1 and 30 °C. On Fig. 1, titanium X-ray maps, obtained by electron probe micro-analysis, illustrate the progressive incorporation of the reactant Fe2Ti into the FeCrW matrix through ballmilling. Each map depicts transversal sections of powder grains where areas enriched in titanium appear bright. After 1 h of ball-milling, Fe2Ti fragments are stuck on the surface of bigger matrix grains. Between 5 and 48 h, those fragments are gradually incorporated inside matrix grains and refined. Since Fe2Ti is fragile and FeCrW is ductile, this is a commonly observed behavior during ball-milling [32]. After 72 h, only a very few micrometric titanium precipitates remain and after 144 h, titanium repartition

Fig. 1. Titanium X-ray maps obtained by EPMA of powder grains resulting of the milling of FeCrW + Fe2Ti + YFe3 + Fe2O3 for different durations of milling at 30 °C under an intensity of 2000 m s2. All images have the same magnification. Top left corner sample is B24 and bottom right corner sample is Ref-B.

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Fig. 3. Scattering cross-sections, measured on powders resulting of the milling of FeCrW + Fe2Ti + YFe3 + Fe2O3 for different durations (from 24 to 144 h) and of FeCrW for 72 h at 30 °C under an intensity of 2000 m s2.

Fig. 2. 3-D APT reconstructions of B72, resulting of a 72 h long milling of FeCrW + Fe2Ti + YFe3 + Fe2O3.

is homogeneous at a micrometer scale. Similar maps are observed for yttrium, meaning that YFe3 follows a similar incorporation into the metallic matrix. Maps for oxygen, and thus the behavior of Fe2O3, are less obvious since oxygen is an element that is difficult to detect by EPMA at low concentrations, as we have here (0.05 wt.%). The powder which underwent 72 h of ball-milling was investigated at an atomic scale by atom probe tomography. The 3-D reconstructions of Fig. 2 show that the evaporated ions Ti, TiO, Y, YO and O appear randomly distributed. This is confirmed by distribution tests. Taking into account the analyzed volume, if clusters are present, their number density is lower than 3.4  1022 m3. In addition, the mean composition of the analyzed volume is Fe–14.96 ± 0.07 Cr–0.44 ± 0.01 W–0.29 ± 0.01 Ti–0.11 ± 0.01 Y–0.17 ± 0.01 O–0.08 ± 0.01 C (at.%). The levels of Y, Ti and O are very close to the nominal values, which reinforces the idea of a homogeneous distribution of these elements, even at nanometer scale. Moreover, the concentrations in Y and O are significantly higher than their solubility limits in Fe at 800 °C (0.029 and 0.06 at.% respectively for Y and O [33,34]). Thus after 72 h of ball-milling, a solid solution oversaturated in Y and O is formed. Finally, powders ball-milled for between 24 and 144 h were analyzed by small angle neutron scattering. In Fig. 3, the scattering cross-sections of those samples plus the matrix sample ball-milled 72 h are plotted. From 24 to 72 h, the scattering intensity decreases. The cross-section of the ODS and non-ODS powders which were both ballmilled for 72 h are equal. Further increasing the ball-milling time results in increased cross-sections. Thus the ODS

powder ball-milled for 72 h contains only nanoporosities, as in the matrix sample. This is in agreement with APT characterization. Nanometric precipitates are present in other samples. Their characteristics are given in Table 3. For B24 and B48, precipitates are likely to be fragments of Fe2Ti, YFe3 and/or Fe2O3, which are progressively disappearing. Indeed, the mean radius and the volume fraction decrease from respectively 1.85 to 1.1 nm and 0.16 to 0.13% between 24 and 48 h of ball-milling. For Ref-B, precipitates with a mean radius of 1.1 nm and a volume fraction of 0.24% must be nanoclusters. Indeed, in Ref. [35], an ODS powder, with a composition Fe–14 Cr– 2 W–1 Ti–0.8 Y–0.2 O (wt.%) was synthesized in identical ball-milling conditions. An APT characterization of the as-milled powder reveals the presence of Cr, Ti, Y and O enriched NC with an average radius of 0.8 ± 0.2 nm and a density of 2.2 ± 0.5  1024 m3. Thus a two-step formation mechanism was observed during ball-milling under those specific milling conditions. First, reactants Fe2Ti, YFe3 and Fe2O3 are dissolved into the metallic matrix until a solid solution, oversaturated in Y and O, is formed. Second, nucleation of nanoclusters begins. 3.2. Nature and form of reactants The following combinations of reactants, in proportions so as to always have identical compositions, were milled: FeCrWTi + Y2O3 nm (Re1); FeCrWTi + Y2O3 (Re2); FeCrWTi + YFe3 + Fe2O3 (Re3); FeCrW + Fe2Ti + YFe3 + Fe2O3 (Ref-B). The first combination is the one commonly used for ODS synthesis [4,17]. The ball-milling conditions were fixed (duration of 144 h, intensity of 2000 m s1 and temperature of 30 °C). The SANS crosssections of the as-milled samples are similar (see Fig. 4).

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Table 3 Characteristics of nanometric precipitates measured by SANS in as-milled and as-annealed ODS powders. The reactants of all those samples were FeCrW + Fe2Ti + YFe3 + Fe2O3. d, I, T, R and f are respectively the duration, the intensity, the temperature, the average radius and the volume fraction. Uncertainties are 5% and 15% respectively for R and f. Name

Ref-B Ref-B-A Ref-B-A2 B24 B24-A B48 B48-A B72 B72-A B100-A T-B T-B-A

Milling conditions

Annealing conditions 2

Precipitates characteristics

d (h)

I (m s )

T (°C)

T (°C)

d (min)

R (nm)

f (%)

144 144 144 24 24 48 48 72 72 100 144 144

2000 2000 2000 2000 2000 2000 2000 2000 2000 2000 2000 2000

30 30 30 30 30 30 30 30 30 30 150 150

– 800 800 – 800 – 800 – 800 800 – 800

– 5 60 – 5 – 5 – 5 5 – 5

1.1 ± 0.06 1.1 ± 0.06 1.1 ± 0.06 1.85 ± 0.09 1.65 ± 0.08 1.1 ± 0.06 1.5 ± 0.08 – 1.5 ± 0.08 1.1 ± 0.06 0.9 ± 0.05 1.5 ± 0.08

0.24 ± 0.04 2.0 ± 0.3 1.8 ± 0.3 0.16 ± 0.02 0.6 ± 0.09 0.13 ± 0.02 1.1 ± 0.2 0 1.4 ± 0.2 1.6 ± 0.2 0.65 ± 0.1 1.5 ± 0.2

milling, the same solid solution should be reached and then the same NC dispersion should be formed, as was observed here. However, it is possible that when at least two of the elements taking part in nanoclusters (Y, Ti or O) are brought with the same reactant, such as Y2O3, the competition between dissolution and nucleation during milling leads to the stationary state without going by the solid solution step. 3.3. Ball-milling intensity and temperature

Fig. 4. Scattering cross-sections, measured on powders resulting of a milling of 144 h with the following reactants: FeCrWTi + Y2O3 nm (ODSRe1), FeCrWTi + Y2O3 (ODS-Re2), FeCrWTi + YFe3 + Fe2O3 (ODSRe3), FeCrW + Fe2Ti + YFe3 + Fe2O3 (ODS-Ref) and FeCrW (M-B).

Moreover, the ODS cross-sections are slightly more intense than that of the matrix sample, which was ball-milled in the same conditions, the extra diffusion being due to NC, as observed in Section 3.1. So, in those specific milling conditions, NC nucleation starts during ball-milling, whatever the reactants nature or form. Consequently, the nature and form of reactants have no influence on nanocluster size and volume fraction and thus their choice could be made on other criteria, such as their cost, their purity or their availability, for example. Here reactants FeCrW + Fe2Ti + YFe3 + Fe2O3 were chosen as the reference state because if needed, concentrations of Y, Ti and O solutes can be set independently from each other. It is very likely that the previously described formation mechanism, deduced from a detailed study with reactants FeCrW + Fe2Ti + YFe3 + Fe2O3, is valid for any reactants. Indeed, if arbitrary reactants are dissolved by

Now we study how ball-milling intensity and temperature influence the kinetics and the stationary state of milling, that is to say the dynamic equilibrium reached after a long enough milling time. To do so, three intensities (1000, 2000 and 4000 m s2) and two temperatures (30 and 150 °C) were tested. The reactants were FeCrW + Fe2Ti + YFe3 + Fe2O3, and the durations were 24, 72 and 144 h. Titanium X-ray maps show that increasing the ball-milling intensity or temperature fastens the reactant dissolution into the matrix (see Figs. 5 and 1 for comparison). Indeed, after 24 h of ball-milling, the density and the size of titanium precipitates in I1-B24 are greater than in B24 and even more than in I3-B24 and T-B24. After 72 h of ball-milling, no titanium precipitates at all are visible for I3-B72 and T-B72, very few are present in ODS-ref and some remain in I1-B72. SANS allows for comparison of the stationary state at the nanometer scale. On Fig. 6, I1-B, I3-B, T-B and RefB scattering cross-sections are plotted. Cross-sections of I1-B, I3-B and Ref-B are exactly superimposed, contrary to T-B. Indeed, the population of precipitates of this latter sample has an average radius of 0.9 nm, comparable to the three others, but a volume fraction of 0.65%, which is much higher. So, in the tested range, the ball-milling intensity has no influence on the milling stationary state. On the contrary, increasing the ball-milling temperature enhances the NC nucleation without inducing growth.

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Fig. 5. Titanium X-ray maps obtained by EPMA of ODS powder grains after a milling with an intensity of 1000 m s2 (I1) or 4000 m s2 (I3) at ambient temperature and with an intensity of 2000 m s2 and a temperature of 150 °C (T). Images have been taken after 24 and 72 h of milling, with the same magnification.

Fig. 6. Scattering cross-sections of ODS powders at the milling stationary state (144 h) with a milling intensity of 1000 m s2 (I1-B), 4000 m s2 (I3B), a milling temperature of 150 °C (T-B) and a milling in the reference conditions, that is intensity of 2000 m s2 and ambient temperature (RefB). Scattering crosssection of a matrix sample (M-B) is also plotted for comparison.

3.4. Annealing Using different ball-milling conditions, ODS powders with different NC dispersions were obtained. Ball-milling is known to be a process with the ability to create a metastable state; however, annealing often brings the system back to equilibrium. The objective here is to determine whether the different ODS powders remain different after

annealing or whether annealing leads all the powders towards the same equilibrium state. To do so, previously as-milled ODS powders were annealed at 800 °C, for durations of 5 min and 60 min, and characterized by SANS. Results are given in Table 3. First, for the reference ODS annealed at 800 °C for 5 min (Ref-B-A), the average radius of the NC dispersion is equal to that after ball-milling (1.1 nm) and the volume fraction has significantly increased (0.24 and 2.0 after ball-milling and after annealing respectively). So a brief annealing has raised the number of NC. The scattering cross-sections of Ref-B-A and Ref-B-A2 (annealed during 60 min) are essentially identical and appear superimposed (see Fig. 7). The same result has been observed on all samples which have been annealed at 800 °C for 1 h. So, it appears that the number and the size of NC do not evolve further after the first 5 min of annealing at 800 °C. This underlines that formation phenomena can be very rapid and that it is of great interest to study the very beginning of these processes. Second, samples milled for different durations all contained dispersions of NC after annealing but with different characteristics: the longer the powder had been milled, the thinner and denser was the NC dispersion. For instance, the average radius and the volume fraction of NC in B24-A were 1.65 nm and 0.6% respectively whereas they are 1.1 nm and 2.0% respectively in Ref-B-A. Samples that had been milled for less than 72 h contained some remaining reactants. It appears that this does not prevent NC nucleation during annealing. The future of the remaining reactants during annealing cannot be predicted with the performed characterizations.

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Consequently, to obtain a fine and dense NC dispersion after a brief annealing, the following as-milled state should be reached: (i) the complete dissolution of reactants, and (ii) the beginning of NC nucleation without being too advanced. This can be obtained through ball-milling with a long enough and/or great enough intensity at ambient temperature, in other words, with a limited overheating. 4. Discussion

Fig. 7. Scattering cross-sections of an ODS powder ball-milled 144 h (RefB), and then annealed at 800 °C for 5 min (Ref-B-A) or 60 min (Ref-BA2). For comparison the cross-section of a matrix sample (M-B) is also plotted.

Third, samples milled in the same conditions of duration, intensity and temperature but with reactants of various forms all exhibit the same SANS cross-sections after annealing. This is also the case for the samples milled for 144 h with three different ball-milling intensities. This was expected since all those powders had equal SANS cross-sections after milling. In other words, powders that are similar after milling have the same NC dispersion after a brief annealing. Finally, powders milled at 150 °C and those milled at ambient temperature are compared after annealing. The average radius and the volume fraction of NC in T-B-A are 1.5 nm and 1.5% respectively. The dispersion of T-BA is broader and less dense than Ref-B-A, although it was denser than Ref-B after ball-milling. Consequently, different as-milled ODS powders remain different after a brief annealing. This seems to be partly related to Y and O solid solution oversaturation in the asmilled powder, which must induce a chemical driving force and thus promote NC nucleation during annealing. Indeed, Ref-B-A, milled for 144 h, exhibits a denser dispersion than B24-A and B48-A. For those samples, the dissolution of reactants was not finished at the end of ball-milling and thus the solid solution was less oversaturated than in Ref-B. Moreover Ref-B-A had thinner NC and a higher NC volume fraction than T-B-A, in which NC nucleation was more advanced than in Ref-B after ball-milling, leading to a less oversaturated solid solution. However, in Ref-B-A, NC dispersion was also denser than in B72-A in which oversaturation after ball-milling should be higher than in Ref-B. In fact, the dissolution of reactants was complete and NC nucleation had not started in B72. Here NC already present in Ref-B seemed to enhance NC nucleation during annealing, although their volume fraction was very low compared to that of the annealed powder (0.24 and 2.0% respectively for Ref-B and Ref-B-A).

Based on the synthesis and characterization of ODS powders obtained through ball-milling of FeCrW + Fe2Ti + YFe3 + Fe2O3 at 30 °C under an intensity of 2000 m s2 and through subsequent annealing, a formation mechanism for nanoclusters can be proposed. Ballmilling starts with some matter exchanges between powder grains, which results in the mixing of reactants by successive failure and welding. This chemical mixing, as defined in Ref. [36], leads all powder grains to equivalent average chemical composition after around 24 to 48 h of ball-milling. At that point, the interiors of grains are not homogeneous at the nanometer scale. The system evolution goes on towards a complete dissolution of reactants at the atomic scale within the metallic matrix and the formation of a solid solution oversaturated in Y and O. As explained by the forced alloy theory [25,37,38], during milling, ballistic jumps are induced by intra particle shearing and disorder the system, while thermally activated jumps tend to bring it toward thermodynamic equilibrium. This can result in metastable states, like an oversaturated solid solution. This first step is finished after 72 h of ball-milling. Then nanocluster nucleation begins. For this to occur, solute species have to diffuse quickly enough in spite of the ambient temperature. The numerous vacancies created by particles shearing make this possible. More precisely, Fu et al. [39] have shown by ab initio calculations in iron that, when vacancies pre-exist, like in a milled system, the O-vacancy pair formation energy essentially vanishes and that an O-vacancy pair attracts atoms that have an affinity for oxygen, such as Y or Ti. Thus this O-vacancy mechanism enables the formation of O-, Ti- and Y-enriched nanoclusters. Afterwards, during annealing at 800 °C, the NC nucleation is greatly enhanced during the first minutes, but suddenly stops without being followed by any coarsening. A qualitative explanation could be that vacancies remaining from milling plus thermal vacancies initially make possible a fast vacancy-based diffusion, leading the formation of many additional NC. However after a few minutes, most of the milling-generated vacancies will have been annihilated, leading to a drop in the vacancy concentration and preventing nucleation and coarsening during the rest of annealing. The two-step ball-milling mechanism, that is, dissolution of reactants and beginning of NC nucleation, is proposed here as a general mechanism for any ODS system. However, when Y and O come from the same reactant,

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Fig. 8. Simplified scheme of the conventional process for ODS steel (a) and the proposed modified process promoting NC formation (b).

such as Y2O3, the stationary state, characterized by the presence of NC, could be reached before dissolving completely the reactants. The kinetics and the stationary state depend on the ball-milling conditions, mainly intensity and temperature. Up to now, ball-milling was believed only to induce the reactants’ dissolution. But most of the time such a proposition was based on characterizations of ODS steel at the end of the process, which obviously does not permit a precise determination of ball-milling’s role [9,12,18]. In the case of as-milled powders characterization, deep nanoscale characterizations are needed since NC nucleation is only just beginning. Recently, some nanoclusters were detected in as-milled ODS powders using atom probe tomography [19,40] and small angle neutron scattering [41], confirming that NC nucleation can be initiated by ball-milling. However Alinger et al. also performed SANS characterization of as-milled ODS powders without detecting any NC [24,42]. Actually this is not in contradiction with the proposed two-step ball-milling mechanism; due to different milling conditions, those as-milled ODS powders could be still in the dissolution step. Here annealing performed at 800 °C was proven to effectively enhance NC nucleation without inducing any coarsening. Alinger et al. [24] and Hoelzer et al. [15] make similar observations. They have also shown that annealing at temperatures of 1100 °C or greater promotes the formation of bigger stoichiometric oxides. Thus it seems favorable to anneal as-milled powders at 800 °C to finish NC nucleation without beginning to coarsen. Then it should be possible to perform consolidation steps at the necessary high temperatures without any coarsening nor dissolution of NC. Once formed, NC are expected to be very stable, as was shown by ageing MA 957 at 1200 °C for 1 h [43]

or at 1300 °C for up to 24 h [8]. Moreover Couvrat et al. observed no evolution of nanocluster size distribution due to hot extrusion [44]. Consequently a modified process is suggested for ODS steels in order to promote NC formation on stoichiometric nano-oxides (see Fig. 8). There are several important differences compared to the conventional process: (i) If the nominal composition is respected, reactants with any phase and granulometry can be chosen; nanometric Y2O3 powders are not needed. (ii) Ball-milling conditions must be controlled; more specifically milling has to be long and/or intense enough, with a limited overheating inside the vial, so as to just induce the beginning of NC nucleation. (iii) An additional annealing of the as-milled powder at 800 °C is performed to finish NC nucleation without inducing coarsening. Any following consolidation steps, like hot extrusion, remain unchanged. It is underlined that to be able to perform consolidation steps the density of NC must not be too high since hardness increases with NC density. So NC density needs to be precisely controlled to obtain a satisfactory compromise between limited hardness and high creep strength. 5. Conclusion Our objective was to study the very beginning of the nano-oxide formation and to deduce from that a method by which the formation of a dense dispersion of nanoclusters is promoted and the coarsening of those nanoclusters is prevented. To achieve this, ODS powders were synthesized in varying ball-milling and annealing conditions and then characterized in the as-milled and as-annealed states by EPMA, SANS and APT.

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The main results are the following: – A two-step mechanism was identified during ball-milling in specific conditions: first, reactants are dissolved into the metallic matrix until an oversaturated solid solution is formed, and second, nanocluster nucleation begins. – There is a dependence on ball-milling parameters: increasing intensity or temperature hastens the dissolution of reactants. Intensity, in the tested range, has no influence on the milling stationary state. On the contrary, when the temperature is increased, the nanocluster density is increased. – Annealing at 800 °C induces a great enhancement of NC nucleation which stops after few minutes without any coarsening occurring. – As-milled ODS powders that are initially different remain different after a brief annealing. To obtain a fine and dense NC dispersion after a brief annealing, dissolution of reactants should have been completed and NC nucleation should have started during ball-milling, but without being too advanced. This can be obtained through ball-milling under relatively long and/or intense enough conditions, with a limited overheating. – A modified process for ODS steel is proposed in order to promote NC nucleation. The main modifications are the control of ball-milling conditions and the annealing of the as-milled powder at around 800 °C before thermomechanical treatment. In the future, a larger range of ball-milling and annealing conditions will be investigated to improve the understanding and the control of nano-oxides formation. Moreover ODS powders will be produced in larger quantities in order to be able to consolidate them and perform mechanical tests. Acknowledgments Authors greatly acknowledge G. Le Cae¨r from the University of Rennes for his interest in this work and for his helpful comments and suggestions. References [1] Klueh RL, Shingledecker JP, Swindeman RW, Hoelzer DT. J Nucl Mater 2005;341:103. [2] Ramar A, Spa¨tig P, Scha¨ublin R. J Nucl Mater 2008;382:210. [3] Sokolov MA, Tanigawa H, Odette GR, Shiba K, Klueh RL. J Nucl Mater 2007;367:68. [4] Ukai S, Mizuta S, Fujiwara M, Okuda T, Kobayashi T. J Nucl Sci Technol 2002;39:778. [5] Sakasegawa H, Chaffron L, Legendre F, Boulanger L, Cozzika T, Brocq M, et al. J Nucl Mater 2009;384:115. [6] Williams CA, Marquis EA, Cerezo A, Smith GDW. J Nucl Mater 2010;400:37. [7] Miller MK, Hoelzer DT, Kenik EA, Russell KF. J Nucl Mater 2004;329–333:338. [8] Miller MK, Hoelzer DT, Kenik EA, Russell KF. Intermetallics 2005;13:387.

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