Precipitation and age-hardening in the Fe–27Co–8Mo alloy

Intermetallics 22 (2012) 33e40

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Precipitation and age-hardening in the Fee27Coe8Mo alloy P. Galimberti, S. Lay*, A. Antoni-Zdziobek SIMaP, Grenoble INP e CNRS - UJF, 1130, rue de la Piscine, BP75, 38402 Saint-Martin d’Hères, France

a r t i c l e i n f o

a b s t r a c t

Article history: Received 10 June 2011 Received in revised form 21 October 2011 Accepted 22 October 2011 Available online 24 November 2011

The principal focus of this work is to identify the microstructural changes occurring in the Fee27Coe8Mo (mass%) alloy on aging. The constitution and microstructure of the alloy have been studied for aging treatments between 800
C and 600
C using X-ray diffraction, transmission electron microscopy and energy dispersive X-ray spectroscopy. For a long aging duration and for the highest temperatures, micrometric precipitates of m phase are found preferentially at grain boundaries of the alloy. At lower temperatures, nanometric precipitates of R phase are observed inside the grains in addition to m submicron precipitates at the grain boundaries. The m precipitates show lattice parameters slightly different from the Fe7Mo6 phase and contain a high density of crystalline defects. The first phases to come out from solid solution on heating the alloy at 600
C are Mo rich precipitates with a cubeecube orientation and ordered domains. These nanometric phases bring a very high hardness to the alloy at the onset of aging. Ó 2011 Elsevier Ltd. All rights reserved.

Keywords: A. Intermetallics A. Ternary alloy systems B. Age-hardening B. Precipitates

1. Introduction The FeeCoeMo alloys have been early recognised to be candidate materials for permanent magnets or for structural applications like tooling [1]. Particularly a high hardness combined with a high ductility can be obtained using appropriate thermal treatments [2,3]. While these alloys are soft after quenching, tempering treatments at moderate temperatures lead to hardness levels in the range of classical high speed steels [4]. The type of precipitates, their size and their distribution influence the hardening behaviour of the alloys. Up to now, most microstructural investigations were performed on FeeMo alloys and few studies have been devoted to the precipitation in FeeCoeMo alloys. In the former studies on FeeMo alloys, investigations were mainly focussed on the first stages of aging. Depending on the temperature, a modulated microstructure, associated with a spinodal decomposition, or the precipitation of plates with a lattice parameter close to Mo, were observed [5e8]. More recently, the presence of Mo rich clusters as small as 2 nm at the beginning of aging was pointed out in these alloys using atom probe analyses [9]. The hardening of the alloys was correlated with the decomposition of the supersaturated matrix [5,8,10]. In addition, two intermetallic phases, the C14 Fe2Mo compound called l and the Fe7Mo6 phase called m, corresponding to different temperature and composition

* Corresponding author. E-mail address: [email protected] (S. Lay). 0966-9795/$ e see front matter Ó 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2011.10.019

domains were examined. These phases belong to the intermetallic phases called tetrahedrally close-packed (TCP) structures that can be described as a stacking of TCP sheets with the Zr4Al3 or MgCu2 structures. Structural defects originating from mistakes in the stacking are very frequent as shown in the Co7W6 compound [11]. Hence, the Fe7Mo6 and Co7Mo6 phases contain a high density of twins and stacking faults parallel or not to the basal plane [12e14]. Moreover, the disordered state of the Fe2Mo and Fe7Mo6 phases was also revealed and interpreted using electron diffraction and high resolution transmission electron microscopy [15e18]. In the ternary FeeCoeMo system, the aging behaviour was studied between 500
C and 650
C for compositions up to Fe43Co21Mo (mass%) alloys [19]. The observations emphasize the presence of a modulated microstructure or reveal Mo rich plate shaped zones like in FeeMo alloys. For the longer aging durations and a low Co content, the l phase appears in the alloys. On the other hand, in the Fe25Co15Mo alloy (mass%), the spinodal decomposition of the matrix during aging at temperatures lower than 475
C was assessed using atom scale and in-situ investigation methods [20,21]. At higher temperatures, a precipitation process takes place and the m phase is the equilibrium phase [22]. The survey of the literature still reveals few available data concerning the constitution, the microstructure and the mechanical properties of intermetallic FeeCoeMo alloys. While negligible hardening is obtained on aging for low Co additions, increasing the Co content increases the hardness and decreases the aging temperature [2]. On the basis of these requirements, the Fe27Co8Mo (mass%) alloy was chosen. This work aims at describing the

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precipitates appearing in this alloy for long annealing treatments at 800, 660 and 600
C to determine the equilibrium phases at these temperatures. Then the sequence of precipitation at 600
C is studied in relation with the hardness behaviour. 2. Experimental procedure The Fe27Co8Mo (mass %) alloy was prepared from mixtures of available commercial powders of high purity (>99.9%) that were cold pressed. To get stable intermetallic phases, specimens were sintered 4 h at 1000
C. They were then solution treated in the onephase field g at 1250
C for seven days in a horizontal furnace under H2 atmosphere then water quenched at room temperature. They were finally aged at 800
C for 11 days and at 660
C for 14 days and subsequently cooled by quenching. For the aging at 600
C, the samples were sintered for 5 days at 1000
C under H2 atmosphere without being previously solution treated, then water quenched. It has been checked using X-ray diffraction (XRD) and transmission electron microscopy (TEM) that no precipitates are present in the as-quenched Fe27Co8Mo alloy. The specimens were then aged for 10, 15, 25, 60 min, 2 h, 24 h, 33 days and 45 days. In all studied alloys, the matrix has the body centered cubic (bcc) crystallographic structure after quenching and aging and will be called a in what follows. The a phase can be the result of a transformation during quenching. Thus the crystallography of this phase is not in any case representative of the phase which can be austenitic at the heat treatment temperature. The experimental approach is supplemented by thermodynamic calculations. Isothermal sections of the FeeCoeMo ternary system between 600
C and 930
C were determined using ThermocalcÔ software (Fig. 1) [23]. The database [24] was constituted by combining the contributions from the binary systems, lacking accurate experimental information for evaluating ternary parameters. The parameters are directly taken from the most recent assessment of the binary FeeCo [25], CoeMo [26] and FeeMo [27e29]. For modelling the m phase, the three sublattice model proposed for the CoeMo system by Davydov et al. [26] has been used, where one sublattice is occupied only by Co and the remaining two are occupied by a substitutional solution. Thus four parameters have been estimated to ensure the compatibility of the description of the m phase in the FeeMo system and to keep the phase diagram of the original assessment using the approach described in [30] (
LMo:Mo:Fe,Mo,
LMo:Mo:Fe,Mo,
LMo:Mo:Fe,Mo taken equal to 360*T and DfGMo:Mo:Fe taken equal to 80*T). The calculations allow predicting the ternary properties from the lower order systems. It is a first approximation to the properties

Fig. 2. Fe30Co15Mo alloy annealed at 930
C for 28 days: bright field image of a micrometer-sized m precipitate surrounded by the Fe rich matrix and diffraction pattern of the grain labelled A viewed along [1,1,0,0]m.

of the ternary system which is useful for identifying the main features of the system. The hardness was determined by micro-Vickers tests to determine the effect of the annealing treatment on the mechanical properties of the alloy. The precipitation sequence was studied by XRD using the D2AM beam line of the European Synchrotron Radiation Facility (ESRF, Grenoble, France). The energy of the incident beam is equal to 19.95 keV (l ¼ 0.0621 nm) what permits to investigate a large volume of the specimens. The microstructure was characterized by scanning electron microscopy in order to get a general view of the phase distribution, by electron probe microanalysis (EPMA), and TEM with a 3010 JEOL microscope equipped with energy dispersive X-ray spectrometry (EDS). TEM samples were prepared by standard ion-milling technique. The electron diffraction patterns were interpreted owing to the software Carine CrystallographyÔ. In this study, the rhombohedral structure of the m phase was described by a hexagonal lattice. The lattice parameters of the Fe7Mo6 phase (a ¼ 0.47546 nm, c ¼ 2.5716 nm) were used [31]. To obtain rather quantitative data for the composition of the submicron precipitates by TEM/EDS, a procedure to determine the CliffLorimer (CL) factors was conducted. These latter relate the ratios of

Fig. 1. Isothermal sections of the FeeCoeMo ternary system (a) at 930
C and (b) at 600
C with the localization of the gross compositions of the studied alloys.

P. Galimberti et al. / Intermetallics 22 (2012) 33e40

35

3. Results and discussion 3.1. Fe27Co8Mo alloy aged at 800
C for 11 days The alloy contains a large number of precipitates about 1 mm in size especially in the grain boundaries of the a matrix (Fig. 3). The precipitates correspond to the m phase and also show some planar defects parallel to the (0,0,0,1) plane. The XRD profile also fits very well with the calculated profile of the Fe7Mo6 phase (Fig. 4). The TEM/EDS analyses of the precipitates give a mean composition of 70Fee28Coe2Mo for the a matrix and 32Fee13Coe55Mo for the m precipitates (mass%). The calculated values are 69.3Fee29Coe1.7Mo and 33.7Fee12.2Coe54.1Mo for the a matrix and m precipitates respectively. A very good fit is obtained between the experimental and calculated values. This composition is then compared with the one of precipitates formed at lower temperatures. Fig. 3. Fe27Co8Mo alloy annealed at 800
C for 11 days: bright field image of a micrometer-sized m precipitate. Diffraction pattern obtained along the [2,1,1,0]m zone axis of precipitate A showing defects parallel to the basal plane.

3.2. Fe27Co8Mo alloy aged at 660
C for 14 days The XRD profile obtained for the alloy annealed at 660
C for 14 days shows peaks less intense and less sharp than after the annealing at 800
C (Fig. 4). The experimental XRD profile does not fit exactly with the Fe7Mo6 compound. The comparison of the profiles after aging at 660
C and 800
C indicates that the position of some peaks is different. While the (1,1,2,0)m peak lies at the same angle, the (1,0,1,10)m and (1,1,2,6)m peaks are shifted towards smaller values. In the assumption of the m lattice, it would indicate a close a parameter and a slightly larger c parameter. The crystallographic structure of the precipitates was studied using electron diffraction. The TEM observation of the alloy indicates the presence of two types of precipitates: some precipitates about 0.2e0.5 mm in size lie in the grain boundaries while smaller precipitates about 50 nm in size are present in the grains (Fig. 5). For the largest precipitates, the EDS analyses give a mean composition of 33Fee13Coe54Mo (mass%) with only a small deviation from this value. The calculated composition of the m precipitates is 32.9Fee10.9Coe56.2Mo. It is quite the same composition than after annealing at 800
C, for precipitates identified as m phase. The large precipitates show a contrast of lines and most associated diffraction patterns exhibit continuous streaks perpendicular to these lines what reveals a high density of crystal defects inside this phase (Fig. 5c). Moreover precipitates usually contain several grains with different orientations (Fig. 6). Electron diffraction experiments were carried out by selecting only one grain using a selection area aperture or by nano-

concentration and the ratios of X-ray intensities for the elements constituting the analyzed phase [32]. The CL factors depend on the TEM/EDS system and on the micro-analysis conditions. They can be determined using a suitable specimen of known composition [33]. For this purpose, the Fe30Co15Mo alloy (mass%) prepared from a mixture of powders of high purity was used. It was melted in argon atmosphere, then was annealed at 930
C for 28 days. Intermetallic precipitates with a size ranging from 0.4 to 3 mm were observed by TEM. Using low-index beam directions, the diffraction patterns could be recognised to correspond to the m phase as expected from thermodynamic calculations (Fig. 1). Each precipitate is often constituted by several grains containing some planar defects, parallel to the (0001) plane (Fig. 2) as observed in m phase [11e14,34]. Quantitative analyses were conducted using EPMA in the matrix and precipitates. A mean composition of 59.1Fee32.2Coe8.7Mo was found for the matrix and 29.7Fee18.6Coe51.7Mo for the precipitates (mass%). The calculated values are respectively 57.2Fee31.3Coe11.5Mo and 30.8Fee16.2Coe53Mo for the matrix and for the precipitates. EDS analyses of the matrix and precipitates were also carried out in the transmission electron microscope and the knowledge of the concentration ratios permitted to determine the CL factors. These data were used to define the composition of the precipitates in the Fe27Co8Mo alloy. Indeed, these latter are too small to be analysed with the EPMA.

b

a

1,0E+05

1,0E+04

αFe

αFe

αFe

Intensity

1,0E+04

1,0E+03

45 days 600°C

x100 45 days-600°C 1,0E+03 x10 14 days-660°C

1,0E+02

μ μ

μ

1,0E+02

1,0E+01

11 days-800°C 1,0E+01

1,1,-2,6 1,0,-1,10 1,1,-2,0 10

15

2θ (°)

20

11 days 800°C

1,1,-2,0μ

1,0,-1,10μ

1,1,-2,6μ

1,0E+00

1,0E+00

5

14 days 660°C

25

30

14

15

16

17

18

2θ (°)

Fig. 4. XRD patterns of the Fe27Co8Mo alloy for several aging treatments. (a) After annealing at 800
C, the diagram reveals the presence of the a matrix and m phase. After annealing at 660
C and 600
C, the position of the (1,1,2,0)m peak remains unchanged unlike the other peaks of the precipitated phase. (b) Magnification of the profiles in the range 2q ¼ 14e18
showing a small shift of the (1,0,1,10)m and (1,1,2,6)m peaks for the precipitated phase.

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P. Galimberti et al. / Intermetallics 22 (2012) 33e40

Fig. 5. Image of the Fe27Co8Mo alloy annealed for 14 days at 660
C. (a) general view, (b) grain boundary precipitate showing a contrast of lines and (c) associated diffraction pattern showing streaks perpendicular to the lines in the precipitate.

beam diffraction. A careful analysis of the diffraction patterns indicates that a large part of the precipitates can be interpreted using the Fe7Mo6 lattice. Moreover, their composition fits well with the m phase. The presence of a great density of defects in the grain boundary precipitates is not surprising owing to the literature on the m phase. Most studies performed in FeeMo alloys, in the stability domain of l or m phases, indicate the difficulty for these phases to get an equilibrium state [15e17]. Some high resolution transmission electron microscopy (HRTEM) investigations have also revealed that the intermetallic precipitates consist of a mixture of nanodomains of various Frank Kasper phases [18,35,36]. m crystals in CoMo, CoW or in Fe base MoCr rich alloys also contain a lot of crystal defects due to the presence of stacking faults and randomly

distributed twin domains [12e14,34]. The simultaneous presence of Co and Fe in the studied alloy could also contribute to the disordering of the m precipitates at low temperature. Diffraction experiments were also carried out inside the grains in order to get crystallographic information on the nanometersized precipitates. The same type of diffraction pattern is always recorded along the <1,1,0> zone axes of the matrix what reveals a special orientation between the a matrix and the precipitates (Fig. 7). The electron patterns could not be interpreted by considering the m phase, nor the l phase for the precipitates. It was found that they fit exactly with the R phase that is also a TCP phase. The R phase has a rhombohedral unit cell that can be described using a hexagonal lattice ( a ¼ 1.0910 nm, c ¼ 1.9354 nm [31]). The diffraction pattern in Fig. 7ced displays the orientation relationship

Fig. 6. Study of a grain boundary precipitate in the Fe27Co8Mo alloy aged for 14 days at 660
C. (a) Bright field image of the precipitate containing several grains called A to D. (bed) Diffraction patterns (DP) related to the grain A in the precipitate. They were recorded using a double tilt specimen holder for several inclinations. Note that the DPs (b) and (d) were recorded by selecting an area comprising the precipitate and the matrix while the DP (c) was obtained using a nano-beam. (e) Stereographic projection showing the three zone axes interpreted using the crystal lattice of the m phase: dark dots refer to poles while open dots correspond to directions. For sake of clarity, the Miller notation was used for this figure.

P. Galimberti et al. / Intermetallics 22 (2012) 33e40

37

Fig. 7. Study of the intragranular precipitates in the Fe27Co8Mo annealed for 14 days at 660
C. (a) Bright field image, (b) electron DP arising from the a matrix viewed along [1,1,0]a (strong spots) and precipitates (weak spots) when two variants of orientation occur for the R precipitates, (c) when only one variant occurs, (d) simulated DP viewed along [1,1,0]a // [5,1,4,0]R for one variant of orientation assuming the same intensity for all the diffracted spots. On the experimental pattern, many reflections are not visible, according to their actual structure factor.

<1,1,0>a // <5,1,4,0>R and (1,1,1)a // (0,0,0,1)R. This is the same mutual orientation than the one found in a Cr rich steel by Dyson et al. [37] that was later confirmed in other steels containing both Mo or W and Cr [38,39]. The occurrence of two variants of this orientation relationship is responsible for the commonly observed diffraction pattern of Fig. 7b. This latter results from the simultaneous orientation relationships <1,1,0>a // <5,1,4,0>R with (1,1,1)a // (0,0,0,1)R and (1,1,1)a // (0,0,0,1)R. Diffraction patterns along <111>a were also recorded in order to confirm the presence of the R phase in the alloy. However, the typical hexagonal pattern corresponding to the <0001>R zone axis, as observed in Fig. 6 of [38] for example, was not recognised. It is likely due to the composition of the precipitates in the present study that influences the intensity of the diffracted beams. Hence, the simulation of the diffraction pattern along <0001>R zone axis, using the atomic positions of [31], indicates that the intensity of most reflections is weak or equal to zero. Despite the lack of the

<0001>R pattern, the R phase is assumed in what follows owing to the exact match of the other results with this compound. The presence of the R phase is surprising in the Fe27Co8Mo alloy annealed for 14 days at 660
C. In the FeeMo system, the R phase is stable at temperatures higher than 1200
C [27] and in the FeeCoeMo system, it is expected to be stable only above 1200
C as well. The diffraction pattern recorded along <1,1,0>a direction for intragranular precipitates is also sometimes observed for grain boundary precipitates. It indicates that some precipitates lying at the grain boundaries in the Fe27Co8Mo alloy are also R phase.

1,0E+08 η-Mo

αFe

μ

αFe

αFe

1,0E+07 45 days

x105

33 days

x10

4

24 h

x10

3

2h

x102

15 min

x10

Intensity

1,0E+06 750

HV30

25 min 1h

2h

15 min

600

+04

1 day

10 min

550

+05

1,0E+03

500

1,0E+02

450

as quenched

400

Aging at 600°C

350 300 0

50

Time (min)

150

24 h

Fig. 8. Hardness of the Fe27Co8Mo alloy versus the annealing time at 600
C.

1,0E+01 5

10

15

2θ (°)

25

30

35

Fig. 9. XRD patterns of the Fe27Co8Mo alloy aged at 600
C for several durations. Compared to the profile of the as quenched alloy, peaks with a small intensity, indicated by an arrow, can be detected after aging for 15 min.

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P. Galimberti et al. / Intermetallics 22 (2012) 33e40

Fig. 10. Microstructure of the Fe27Co8Mo alloy aged for 15 min at 600
C. (a) Bright field image free of intermetallic precipitates. (b) Dark field (DF) image using the (1,0,0)a diffracted spot and (c) corresponding DP with a zone axis close to [0,0,1]a showing the presence of the superlattice (1,0,0)a and (0,1,0)a spots. (d) DF image of the h-Mo platelets using the (0,2,0)h reflection (e) or the (2,0,0)h reflection and (f) corresponding DP.

3.3. Fe27Co8Mo alloy aged at 600
C The hardness of the alloy was measured after different annealing durations at 600
C (Fig. 8). The curve shows a significant increase at the beginning of the aging treatment. The constitution of the alloy was investigated by XRD in order to determine the crystallographic structure of the precipitates responsible for the hardening behaviour. In addition to the a matrix, the XRD profiles indicate for the shortest annealing durations, especially after 2 h, the presence of a phase with a bcc structure and a lattice parameter of about 0.310 nm (Fig. 9). It is close to the lattice parameter of Mo (0.314 nm). These precipitates likely correspond to the Mo rich h metastable phase, already described in the literature for FeeMo alloys [5,9]. After 33 days and 45 days, the XRD profile is close to the one measured after 14 days at 660
C : peaks related to the precipitated phase do not fit exactly with the Fe7Mo6 lattice (Fig. 4). The microstructure of the alloy annealed for 15 min was observed by TEM. At this early stage of aging, no intermetallic precipitates are visible on the bright field images (Fig. 10a). Firstly,

superlattice (1,0,0) reflections can be observed on the diffraction patterns of the a matrix and dark field experiments permit to image in some areas of the sample small domains about 100 nm in size (Fig. 10b, c). This observation suggests the presence of ordered domains in the matrix after a small aging time. The ordering is not confirmed on the XRD profile where no (1,0,0)a peak, expected at 14.4
, is detected (Fig. 9). It is to note that such domains were also revealed from high resolution transmission electron images in the Fe20Co18W alloy annealed at 800
C for a short time but not detected on the XRD profiles [40]. The formation of these ordered domains that would be a consequence of Co addition should affect the hardness behaviour of the alloy. Some striation is also visible on the TEM images under some diffracting conditions and can be related to the presence of Mo rich clusters. In order to image the h phase, the matrix was oriented along <001>a. Due to the similarity of a and h lattices and the rather close lattice parameters (15% of misfit), the two phases are expected to have the same orientation. Then dark field images were recorded by positioning the objective aperture in such a way to

Fig. 11. Microstructure of the Fe27Co8Mo alloy annealed at 600
C. (a) Bright field image of the alloy aged for 15 min and (b) example of DP recorded after 15 min of aging along <1,1,0>a zone axis with reflections shown by arrows that do not arise from the matrix. (c) Bright field image of the alloy aged for 45 days showing nanometric precipitates in the grains and sub-micron ones in the grain boundaries (GBP).

P. Galimberti et al. / Intermetallics 22 (2012) 33e40

include the {2,0,0}h or {0,2,0}h reflection. Thin platelets about 1.5 nm in thickness and 8 nm in length parallel to the (1,0,0)a and (0,1,0)a planes appear on the micrographs (Fig. 10def). Regarding the literature on the early stages of precipitation in FeeMo alloys, the presence of Mo rich platelets is not surprising. They are expected to appear first with a facetted spherical shape, then to evolve towards an elongated shape when they coarsen and the concentration of Mo atoms can reach about 96 at% in these precipitates [9]. Moreover, h platelets with a size ten times larger than in this study were observed in the Fe27Mo alloy (mass%) aged at 600
C for 15 min [8]. Plate-shaped h zones were also observed in FeeCoeMo alloys [19]. However, the h phase does not seem to be a mandatory step before the precipitation of the stable intermetallic phase. In the Fe27Co15Mo alloy (mass%), no h phase was detected during the heating of the alloy before the precipitation of the m phase [22]. In addition to the superlattice reflections, other reflections with a low intensity, different from the a matrix, are present on the diffraction patterns. It suggests that intermetallic nanosized precipitates begin to form in the alloy, although they are not readily observed by TEM (Fig. 11a, b). When aging proceeds, the h phase disappears and intermetallic phases develop (Fig. 9). The microstructure of the alloy after annealing 45 days at 600
C is very similar to the one appearing after 14 days at 660
C except the smaller size of the precipitates, close to 200 nm in the grain boundaries and 20e50 nm inside the grains (Fig. 11c). The electron diffraction patterns arising from the grain boundary precipitates agree with the m phase as expected from thermodynamic data (Fig. 1) and those associated with the intragranular precipitates are similar to the ones recorded at 660
C, so they correspond to the R phase (Fig. 7). Moreover, the mean composition of the grain boundary precipitates is 36Fee10Coe54Mo (mass%), very close to the composition at higher temperature and close to the calculated equilibrium composition, 32.7Fee10.4Coe56.9Mo, for the m precipitates. The composition of the intragranular precipitates could not be determined because of their small size. The position of the experimental peaks on the XRD patterns after 45 days was compared with those recorded from the R phase in the study of Sinha et al. [31]. It is observed that they do not fit with the R phase. This result is not surprising: from previous work in the FeeCoeW system, it was observed that a signal arising from nanometric intermetallic precipitates was barely detected by XRD [40]. The peaks of the precipitated phase are likely provided by the intergranular precipitates. They fit with the m phase but with lattice parameters slightly different from Fe7Mo6. As an indication of the deviation from the equilibrium m phase, the lattice parameters of this phase are calculated from the XRD profile assuming the same a value as for Fe7Mo6 (a ¼ 0.47546 nm) since no shift of the (1,1,2,0) peak is depicted. The c parameter is found equal to 2.63  0.02 nm after 45 days instead of 2.5713 nm for Fe7Mo6. Owing to the measured composition of the precipitates, the change in lattice parameters cannot be related to a change in Co or Mo content. Such a small deviation of the lattice parameters was already observed for the l phase in the Fe20Co18W alloy [41]. On the basis of the literature on l or m phases, the slightly different lattice parameters, already noticed at 660
C, can be attributed to the disordered state of the crystals when annealing at a low temperature. The examination of the microstructure reveals therefore that the strong increase in hardness at the beginning of the aging is related to nanometric metastable phases. It was checked that the R phase is not present after quenching the alloy from 1250
C. The R precipitates form during the aging of the alloy at low temperature inside the grains while the stable m phase forms at grain boundaries. This result suggests that the

39

formation of the equilibrium m phase is related to the enhanced atomic diffusion at the vicinity of the grain boundaries. In the ferrite grains, the precipitation of the intermetallic R phase could be due to the structural proximity between the cells of the a matrix and the R phase. Hence, the transformation involves deformations less than 3% and atomic movements smaller than inter-atomic distances [37,38]. 4. Conclusions The microstructural investigations in the Fe27Co8Mo alloy annealed between 800
C and 600
C provide new information on the precipitation behaviour in FeeCoeMo alloys. - For a long aging duration at 800
C, the stable intermetallic m phase precipitates with very few defects and the crystal lattice fits exactly with the Fe7Mo6 lattice. The composition of the precipitates agrees with the one calculated from available thermodynamic data. After a long aging time at 600
C and 660
C, precipitates containing a high density of defects form in the grain boundaries with a composition and a lattice close from the m phase. Inside the grains, nanometric precipitates of R phase have been found. The low annealing temperatures do not give rise to the equilibrium structure even after a long tempering period. - The following precipitation sequence was identified when aging the alloy at 600
C: nanometric Mo rich platelets, submicron ordered domains and possibly nanometric intermetallic precipitates form in the first stages of aging. They are responsible for the marked increase of the alloy hardness at the onset of aging. On further aging, Mo rich platelets and ordered domains disappear, nanometric particles of R phase and m precipitates are found in the ferrite grains and at the grain boundaries respectively. The a / R transformation requires only small atomic movements and give rise to a limited strain [37,38] what explains the preferential precipitation of R phase in the grains where diffusion is more limited than in grain boundaries. - Further HRTEM observations and atom probe analyses would be helpful to characterise the structure and composition of the nanometric phases appearing during aging of this alloy. Acknowledgements The authors thank Marc de Boissieu and Françoise Bley for their assistance at the D2AM synchrotron beamline (ESRF, Grenoble, France). References [1] Köster W. Mechanical and magnetic precipitation hardening of the ironecobaltetungsten and ironecobaltemolybdenum alloys. Arch Eisenhüttenwesen 1932;5:17e23. [2] Edneral AF, Zhukov OP, Perkas MD. Effect of cobalt on aging of martensite and ferrite in FeeCoeW and FeeCoeMo alloys. Sci Heat Treatment 1974;16: 840e3. [3] Karpov MI, Wnukov VI, Medved NV, Danninger H. PM technology for tool steels with intermetallic hardening. In: Proc. Powder Metall. World Congress, Granada; 1998. High Alloy Steels, 519e524. [4] Danninger H, Harold CH, Gierl CH, Ponemayr H, Daxelmueller M, Simancik F, et al. Powder metallurgy manufacturing of carbon-free precipitation hardened high speed steels. Acta Physica Polonica A 2010;117:825e30. [5] Hornbogen E. Demixing of ironemolybdenum and ironetungsten mixed crystals. Zeitschrift für Metallkunde 1961;52:47e56. [6] Higgins J, Wilkes P. Precipitation in the FeeMo and FeeAu systems. Phil Mag 1972;25:599e623. [7] Takagishi S, Miyazaki T, Mori H, Kozakai T. Modulated structure in ironemolybdenum alloys. Scripta Met 1979;13:553 (whose figures are in 1979;13:912 as erratum).

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[8] Miyazaki T, Takagishi S, Mori H, Kozakai T. The phase decomposition of ironemolybdenum binary alloys by spinodal mechanism. Acta Metall 1980; 28:1143e53. [9] Isheim D. Metastable phase formation during the decomposition of Fe-20 at.% Mo. Acta Materialia 2000;48:2873e83. [10] Ustinovshikov Y, Shiren C, Shirobokova M. Laves phase formation in solids. J Mat Sci 1994;29:1411e6. [11] Carvalho PA, De Hosson JThM. Stacking faults in the Co7W6 isomorph of the m phase. Scripta Mater 2001;45:333e40. [12] Zhu J, Ye HQ. On the microstructure and its diffraction anomaly of the m phase in superalloys. Scripta Metall Mater 1990;24:1861e6. [13] Stenberg L, Andersson S. Electron microscope studies on a quenched FeeMo alloy. J Sol Stat Chem 1979;28:269e77. [14] Hiraga K, Yamamoto T, Hirabayashi M. Intermetallic compounds of the m- and P-phases of Co7Mo6 studied by 1 MV electron microscopy. Trans Jap Inst Metals 1983;24:421e8. [15] Michel JP. Etude de la précipitation dans l’alliage fer-6% atomique molybdène. Mém Sci Rev Métallurg, LXVIII 1971;11:785e92. [16] Hirata A, Koyama Y. Disordered atomic column state in FeeMo alloys. Phys Rev B Condensed Matter Mater Phys 2004;70:134203e7. [17] Hirata A, Iwai A, Koyama Y. Characteristic features of the Fe7Mo6-type structure in a transition-metal alloy examined using transmission electron microscopy. Phys Rev B 2006;74:1e4. 054204. [18] Carron D, Chemelle P, Michel D, Hytch MJ, Portier R. Icosahedral related precipitation at a nanometer scale in a superferritic steel. J Non-Crystal Solids 1993;153, 154:473e7. [19] Kosakai T, Aihara H, Doi M, Miyazaki T. Spinodal decomposition in FeeMoeCo and FeeMoeV ternary systems. Trans ISIJ 1985;25:159e64. [20] Eidenberger E, Schober M, Schmoelzer T, Stergar E, Staron P, Leitner H, et al. Analysis of the multistage phase separation reaction in Fe-25 at%Co-9 at%Mo. Phys Stat Sol A 2010;10:2238e46. [21] Eidenberger E, Schober M, Staron P, Caliskanoglu D, Leitner H, Clemens H. Spinodal decomposition in Fe-25 at%Co-9 at%Mo. Intermetallics 2010;18: 2128e35. [22] Leitner H, Schober M, Clemens H, Caliskanoglu D, Danoix F. Precipitation behaviour of an FeeCoeMo-alloy during non-isothermal ageing. Int J Mat Res 2008;99:367e74. [23] Sundman B, Jansson B, Andersson JD. The ThermoCalc databank system. Calphad 1985;9:153e90. [24] Antoni-Zdziobek A. FeCoMo database, unpublished revision.

[25] Guillermet AF. Critical evaluation of the thermodynamic properties of the ironecobalt system. High Temp High Press 1988;19:477e99. [26] Davydov AV, Kattner UR. Revised thermodynamic description for the CoeMo system. J Phase Equilibria 2003;24:209e11. [27] Guillermet AF. An assessment of the FeeMo system. Calphad 1982;6:127e40. [28] Andersson JO. A thermodynamic evaluation of the FeeMoeC system. Calphad 1988;12(1):9e23. [29] Frisk K. An experimental and theoretical study of the phase equilibria in the FeeMoeNi system. Met Trans A 1992;23(2):639e49. [30] Saunders N, Miodownik AP. Calphad, calculation of phase diagrams. A comprehensive study. In: Cahn RW, editor. Pergamon materials series. Pergamon; 1998. [31] Sinha AK, Buckley RA, Hume-Rothery W. Equilibrium diagram of the ironemolybdenum system. J Iron Steel Inst 1967;205:191e5. [32] Cliff G, Lorimer GW. The quantitative analysis of thin specimens. J Microsc 1975;103:203e7. [33] Williams DB, Carter CB. Transmission electron microscopy e spectrometry IV. New York: Plenum Press; 1996. [34] Carvalho PA, Haarsma HSD, Kooi BJ, Bronsveld PM, De Hosson JThM. HRTEM study of Co7W6 and its typical defect structure. Acta Mater 2000;48: 2703e12. [35] Kuo KH, Ye HQ, Li DX. Tetrahedrally close-packed phases in superalloys: new phases and domain structures observed by high-resolution electron microscopy. J Mater Sci 1986;21:2597e622. [36] Ye HQ, Li DX, Kuo KH. Fivefold symmetry in real and reciprocal spaces. Ultramicroscopy 1985;16:273e8. [37] Dyson DJ, Keown SR. A study of precipitation in a 12% CreCoeMo steel. Acta Met 1969;17:1095e107. [38] Redjaïmia A, Morniroli JP, Donnadieu P, Metauer G. Microstructural and analytical study of heavily faulted Frank-Kasper R-phase precipitates in the ferrite of a duplex stainless steel. J Mater Sci 2002;37:4079e91. [39] Yamamoto K, Kimura Y, Mishima Y. Effect of matrix microstructure on precipitation of Laves phase in Fee10Cre1.4W(eCo) alloys. Intermetallics 2006;14:515e20. [40] Galimberti P, Lay S, Antoni-Zdziobek A, Coindeau S, Véron M, Bley F, et al. Precipitation and hardening behaviour of the Fe20Co18W alloy aged at 800
C. Intermetallics 2011;19:556e62. [41] Galimberti P, Lay S, Antoni-Zdziobek A. Laves phase formation in the Fe20Co18W alloy annealed at 800
C. Solid State Phenomena 2011;172e174: 517e24.