Microstructural evolution of Alloy 709 during aging

Materials Characterization 154 (2019) 400–423

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Microstructural evolution of Alloy 709 during aging a,⁎

a

a

a

b

Rengen Ding , Jin Yan , Hangyue Li , Suyang Yu , Afsaneh Rabiei , Paul Bowen a b

a

T

School of Metallurgy and Materials, University of Birmingham, Edgbaston, Birmingham B15 2TT, UK Department of Mechanical and Aerospace Engineering, North Carolina State University, 911 Oval Dr, Raleigh, NC 27695-7910, USA

ARTICLE INFO

ABSTRACT

Keywords: Alloy 709 Microstructure Aging TEM

The creep-resistant austenitic stainless steel Alloy 709 (Fe-20Cr-25Ni (wt%) based steel) is being investigated as a candidate structural material for the next generation fast neutron reactors at service temperature of 500–550 °C. However, the study of microstructural evolution of Alloy 709 during aging is lacking. In this study, thus, the microstructure of Alloy 709 has been investigated using electron microscopy in the as-received state and after static aging at 550, 650 and 750 °C. The results show that the prominent precipitate in the as-received Alloy 709 is Nb(CN), with the rod-like Z phase (CrNbN) observed very occasionally. After aging at 550 °C even up to 2000 h, no significant microstructure change was observed, which means that Alloy 709 is fairly stable at 550 °C. Aging at 650 °C produced globular M23C6 phase on grain boundaries, plate-like M23C6 carbides at twin boundaries and in the grain interior, and blocky M23C6 carbide nucleated on Nb(CN). Fine dispersoid Z phases were found on dislocations after aging at 650 °C for 500 h; their amount increases with aging time and temperature. (Cr,Mo)3(Ni,Fe)2SiN θ phase forms at grain boundaries after aging at 650 °C for 1000 h. After aging at 750 °C, θ phase nucleated on M23C6 carbide and a transformation of M23C6 to θ phase was found, which suggests that θ phase is the more stable. Aging at 550 °C promotes the segregation of Cr and Mo to grain boundaries whereas clear Cr depletion was observed at 650 °C due to the precipitation of Cr-rich M23C6 carbides at grain boundaries. Such depletion nearly disappears at 750 °C. The implications of aging for the subsequent mechanical behaviour of Alloy 709 are discussed briefly.

1. Introduction Austenitic stainless steels are the most widely used stainless steels across a wide variety of aerospace, powerplant and chemical applications, due to their excellent corrosion resistance and mechanical properties at high temperatures. To meet continuously increasing service temperature, some advanced heat-resistant austenitic stainless steels with complex chemical compositions, such as HR3C (Sumitomo Metal, Japan), Sanicro 25 (Sandvik, Sweden), DMV 310 N (Salzgitter, Mannesmann, Germany) and NF 709/Alloy 709 (Nippon Steel, Japan) have been developed. Among these, NF709 was developed from Nb/Ti stabilised 20Cr25Ni stainless steel. The relatively newly developed Alloy 709 has a similar chemical composition to NF709 but without Ti, thus having no coarse TiN. Due to its superior high temperature performance and good sodium compatibility, Alloy 709 has been investigated as a candidate structural material for sodium-cooled fast reactors [1] and other forms of advanced nuclear reactors. Unlike conventional solid solution austenitic stainless steels, these advanced austenitic stainless steels develop a variety of precipitates during aging at different temperatures. Knowles [2] reviewed a number ⁎

of microstructural studies on Nb stabilised 20Cr25Ni steels, and established an approach to the prediction of precipitates in terms of alloy composition. He suggested that aging of alloy compositions with stabilisation ratios [wt%Nb/wt%(C + N)] below 7.7 should lead to the precipitation of NbC and M23C6 while aging of alloy compositions with a ratio > 7.7 should result in the formation of NbC and M6C. Additionally, G-phase (Ni16Nb6Si7) was reported in Nb-stabilised 20Cr25Ni steel during creep at 750 and 850 °C, and was believed to form from pre-existing NbC particles as a result of segregation of Si to them [3]. The transformation of NbC to G-phase in the same type of steel was also confirmed even at 500 °C [4]. Naoi et al. [5] found that aging at 700 °C produced a large number of carbides at grain boundaries and adjacent areas in NF709, leading to a decrease in Charpy impact toughness to 30-40 J after aging for 1000 h. Furthermore, Sourmail and Bhadeshia [6] investigated the precipitation behaviour of NF709 and a NF709 variant at 750 and 800 °C, and found that M23C6 (after a short age) and Cr3Ni2SiX(X: mainly N) (after a long age) precipitated at grain boundaries and twin boundaries, together with the formation of Z-phase (CrNbN), predominantly on dislocations. They also found that despite the NF709 variant having a similar chemical

Corresponding author. E-mail address: [email protected] (R. Ding).

https://doi.org/10.1016/j.matchar.2019.06.018 Received 15 April 2019; Received in revised form 11 June 2019; Accepted 11 June 2019 Available online 16 June 2019 1044-5803/ © 2019 Elsevier Inc. All rights reserved.

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Table 1 Chemical composition of Alloy 709. Element Wt%

Cr

Ni

Mn

Mo

Si

Nb

Ti

N

C

S

P

19.69

25.0

0.88

1.46

0.28

0.23

< 0.01

0.14

0.063

< 0.001

< 0.05

Dark – >15° Red – twin Green – <5°

(b)

(a) 60 50

Frequency, %

40 30 20 10 0 0

10

20

30

40

Misorientation angle,

50

60

o

(c) Fig. 1. (a) Inverse Pole Figure (IPF) EBSD map of as-received Alloy 709, (b) EBSD grain boundary map showing high angle grain boundaries in black, Σ3 twin boundaries in red and low angle grain boundaries in green and (c) a profile of misorientation angles showing ~50% Σ3 (twin) boundaries. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

(b)

(a)

Fig. 2. (a) A typical BSE image of as-received Alloy 709 showing a band-like distribution of white particles, (b) the EDS spectrum shows that those white particles are Nb-rich carbonitride. Note: RD – rolling direction.

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(a)

(b)

(c)

(d)

4000

Nb

Cr Fe

Counts

3000

Cr N 2000

Cr Ni

Fe

1000

Fe 0 0

1

2

3

(f)

(e)

4 keV

5

6

7

8

(g)

Fig. 3. (a) Bright field (BF) scanning transmission electron microscopy (STEM) image and (b) high angle annular dark field (HAADF) STEM image of the as-received material show nano-sized precipitates. (c) EDS mapping suggests that these particles are Nb carbonitrides, (d) selected area diffraction patterns (SAD) recorded along [111] from a Nb(CN) precipitate, (e) HAADF –STEM image of rod shaped Z-phase, (f) EDS spectrum from the Z-phase and (g) SAD pattern taken along [010] zone axis from the Z-phase.

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(a)

(b)

60

Fe Ni Cr Mo

wt%

40

20

0 0

20

40

60

nm

(c) Fig. 4. BF-STEM images taken from the samples aged at 550 °C for (a) 500 h and (b) 2000 h, showing that aging did not cause the formation of other particles beyond the already present Nb(CN). (c) EDS linescan across a grain boundary revealing that aging promotes the segregation of Cr and Mo to the grain boundary.

composition to NF709, it exhibits different precipitation behaviour. These studies have shown clearly that even though the matrix has the same composition, slight variations of minor elements can lead to very different precipitation scenarios. For instance, G phase was found to be dominant in some cases whereas another Si-containing phase (Cr3Ni2SiN) was observed in the 20Cr25Ni-based steels with high nitrogen such as NF709. This seems to be a result of nitrogen stabilising the Si-containing phase but more evidence is required. This is one of aims of this study. The above studies also demonstrated the formation of a large amount of Cr-rich carbide (e.g. M23C6) at grain boundaries during aging, which could lead to Cr depletion at the grain boundaries, thus degrading corrosion resistance. However, no relative study has been

reported for NF709 and Alloy 709. The microstructural evolution during aging at elevated temperature is of considerable importance, both to identify any precipitates that might not occur during short-term heat treatments and to ensure that no phase forms that is detrimental to creep properties and that could invalidate the extrapolation of shortterm creep test data. There have been some studies of microstructural evolution in NF709, but there is a lack of investigation of Alloy 709. Thus, this paper aims to fill this gap by reporting grain boundary chemistry and precipitate identification in aged Alloy 709 samples. This study also provides a data base for Alloy 709 which will be used as a structural material for sodium-cooled fast reactors.

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(a)

(b)

(c)

(d)

Fig. 5. (a) BF and (b) HAADF STEM images taken from the sample aged at 650 °C for 500 h, showing the nearly continuous films with some globular particles on the grain boundaries and a large number of small globular particles and plate shaped particles in the grain interiors. (c) Higher magnification image of the GBs in Fig. 5a and (d) higher magnification image of the twin in Fig. 5a illustrating that plate shaped particles formed particularly at the incoherent twin boundary. The arrowed particle is analysed in Fig. 6.

2. Experimental procedure

is given in Table 1. Samples were sealed in quartz tubes which were filled with argon and aged in a furnace at 550, 650 and 750 °C for 500, 1000, 2000 h, respectively. Since Alloy 709 is regarded as a candidate structural material for the sodium-cooled fast reactor where its service temperature is 500–550 °C, the study of its microstructural evolution at 550 °C is of especial importance. An aging time of 2000 h is still short compared

The Alloy 709 used in this study was fabricated using vacuum-induction melting and electro-slag remelting by Carpenter Technologies. The ingot was homogenized at 1250 °C for 4 h, hot forged at 1100 °C and then rolled at 1100 °C. The hot-rolled plate was finally annealed at 1100 °C for 2 h, followed by water-quenching. Its chemical composition

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(a)

(c)

(b)

(d)

Fig. 6. (a) HAADF-STEM image of the sample aged at 650 °C for 500 h and the corresponding EDS maps showing that the GB particles are (Cr,Mo)-rich carbides with some fine Nb-rich phases, while the particles in the grain interior include Nb(CN), composite particles with a Nb(CN) core and a shell of (Cr,Mo)-rich carbide, and plate shaped (Cr,Mo)-rich carbide. SADs were recorded along (b) [101], (c) [121] and (d) [011] from the M23C6 carbide (arrowed in Fig. 5c), showing that this GB's M23C6 carbide has a cube to cube orientation relationship with matrix at upper side of the GB. Table 2 Composition of (Cr,Mo)23C6 in Alloy 709 (at.%). Temperature, °C

Time, h

650

500 1000 2000 500 1000 2000

750

Cr 65.9 66.9 68.5 73.5 74.1 74.9

± ± ± ± ± ±

Mo 1.0 2.4 2.7 2.5 1.9 2.2

4.7 5.1 5.2 5.4 5.9 6.3

to its service life and thus higher temperatures (650 and 750 °C) were also used, to promote microstructural evolution. As the precipitation behaviour of NF709 at 750 °C has been reported, the study of Alloy 709 at 750 °C will reveal the effect of minor element variation (i.e. Ti) on

± ± ± ± ± ±

Ni 0.3 0.3 0.4 0.5 0.7 0.7

5.7 5.3 4.8 4.4 4.6 4.6

± ± ± ± ± ±

Fe 0.8 1.2 1.0 0.4 0.8 1.3

22.0 21.0 19.7 15.3 13.5 12.6

± ± ± ± ± ±

1.6 2.3 2.3 2.5 2.3 1.5

precipitate behaviour. Specimens for optical and scanning electron microscopy (SEM) were prepared following a standard metallographic procedure and etched with a solution of 10 pct (by mass) oxalic acid in distilled water. The

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(a)

(b)

(c)

(d) (caption on next page)

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Fig. 7. (a) A HAADF-STEM image illustrating that Nb(CN) acts as a nucleation site for other particles and EDS maps showing those particles are (Cr,Mo)-rich carbides. (b) BF-TEM image of the particle on the left side of the HAADF image showing the formation of (Cr,Mo)-rich carbides on the surface of Nb(CN), (c) high resolution TEM image of part of the composite particle and (d) composite SADs from the MX - Nb(CN), M23C6 and matrix in Fig. 7b revealing the orientation relationship between those three phases. Note: the sample drift caused distortion of the EDS maps.

70 60 50

Fe Cr Ni C Mo

at.%

40 30 20 10 0 0

10

20

30

40

50

nm

(b)

(a)

(c)

(d)

Fig. 8. (a) A BF-STEM image taken from the sample aged at 650 °C for 500 h showing globular and plate shaped particles, (b) EDS linescan across a plate revealing that the particle is (Cr,Mo)-rich carbide, (c) SAD taken along [112] from the matrix and the particle confirming that it is M23C6 and (d) high magnification image of the dislocation (arrowed in Fig. 8a) showing Z-phases, Nb(CN) and (CrMo)23C6 on the dislocation.

grain size was measured using a Oxford Instruments Nordlys EBSD (electron backscatter diffraction) detector with a step size of 2 μm. Thin foils 3 mm in diameter for transmission electron microscopy (TEM) were electropolished in a solution of 10% perchloric acid + 90% ethanol at 20 V and −20 °C, using a twin-jet electropolisher. TEM observations were carried out on FEI Tecnai F20 and Talos F200X microscopes operating at 200 kV, both equipped with a Silicon Drift

Detector (SDD) for energy dispersive X-ray spectrometry (EDS). 3. Results and analysis 3.1. Microstructure of as-received sample The average grain size of the as-received material is 45 μm, which

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Twin

Matrix

(a)

(b)

Matrix M23C6

at matrix side

(c)

(d)

Twin M23C6

at twin side

(e) Fig. 9. (a) A HAADF-STEM image taken from the plate shaped particles at the incoherent twin boundary in Fig. 5d and the corresponding EDS maps. (b) [011] SAD from the matrix and the particles in the matrix, (c) [011] SAD from the twin and the particles in the twin, (d) [011] SAD from the boundary and particles on both sides and (e) schematic diagram of Fig. 9d showing the orientation relationship between the carbides in the twin and the matrix, and between the carbides in the matrix and the twin. Note: small diffraction spots are from the carbides, in Fig. 9d, ts – twin side, ms – matrix side.

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(a)

(b)

(c)

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Fig. 10. BF-STEM images of the sample aged at 650 °C for 1000 h show (a) a nearly continuous distribution of the particles including Si-rich θ phase at the grain boundaries and (b) the formation of plate shaped particles at the ITB but with no evidence of the particles on the CTB, (c) EDS maps for the marked area in Fig. 10a revealing that most of the particles at GBs are M23C6 with a new Si-rich θ phase. (d) More Z-phase particles nucleated on dislocations compared with the sample aged at 650 °C for 500 h (d).

Mo and C (Fig. 6a). The spot EDS analysis reveals that the particle has ~66Cr, ~22Fe and ~5Mo (at. %) (Table 2). SAD patterns taken from a grain boundary particle along different zone axes are shown in Figs. 6bd. The analysis of diffraction combined with EDS indicates that the grain boundary particle is an fcc structure M23C6 carbide with lattice parameter a = 1.06 nm and having a cube-to-cube orientation relationship with the matrix at one side of the grain boundary. Some of the Nb(CN) particles were wrapped partially or completely by (Cr,Mo)rich carbides (Figs. 6 and 7). High resolution (HR) TEM and SAD confirm that the (Cr,Mo)-rich carbide is M23C6, and indicate that the Nb (CN) and M23C6 have an orientation relationship with the γ matrix: [001]Nb(CN)//[101]M23C6//[101]γ (Fig. 7). The 500 h aging also produced some plate shaped particles (Figs. 5 and 8), demonstrated via EDS (Figs. 6 and 8) and SAD (Fig. 8) to be M23C6 carbide. It should be mentioned here that very occasionally, some fine particles were found on dislocations (Fig. 8d) as well as Nb(CN) and M23C6. EDS analysis suggested that they are a (Nb,Cr)-rich nitride, which is proposed to be Z-phase. An interesting observation is that some platelike particles formed on incoherent twin boundaries (ITBs) but not on coherent twin boundaries (CTBs) (Figs. 5d and 9a). EDS and SAD suggest that they are M23C6 (Fig. 9), with some Nb-rich particles (probably Nb(CN)) (Fig. 9). Additionally, SAD analysis also indicates that the plates in the matrix keep orientation relationship (OR) with the twin whereas the plates in twin have OR with the matrix. This longer aging did not clearly change the morphology of the GB particles, but EDS mapping reveals that a new Si-containing phase formed there (Fig. 10c). The Si-rich phase is proposed to be θ phase; its detailed identification will be given in section 3.2.3. The 1000 h aging produced more Z phase predominantly on dislocations (Fig. 10) but no evidence for particles at CTBs (Fig. 10b). Further aging (2000 h) did not lead to the formation of new phases at the grain boundaries (Fig. 11) but produced more Si-containing θ phase. Since M23C6 carbide is rich in Cr while θ phase contains Si, the superposition of Si and Cr-EDS maps reveals that some θ phases grew from the globular M23C6 carbides (white arrow) and from the end of the M23C6 plates (black arrow) (Fig. 11). Additionally plate M23C6 formed at coherent twin boundaries (Fig. 12). EDS and HR-TEM (including diffraction patterns (not shown here)) confirm that the plate-like phase is still M23C6. It should be mentioned here that very occasionally the very tiny particles of Z phase were attached to the M23C6 plates (Fig. 12d).

was determined using EBSD, as shown in Fig. 1. ~ 50% of boundaries are Σ3 twin boundaries. No evidence of strong texture was found in the as-received samples, which means that annealing at 1100 °C alleviated the texture produced by rolling. A typical backscattered electron (BSE) image (Fig. 2(a)) shows a band-like distribution of white particles. EDS analysis indicates that these white particles are Nb-rich carbonitride (i.e. Nb(CN)) (Fig. 2(b)). Such large Nb(CN) particles could form during casting and then break down and align along the rolling direction during rolling. Thus, this type of carbonitride with size in the range 1–10 μm is designated primary Nb(CN). A considerable number of nano-sized particles is also found, as shown in Fig. 3. EDS mapping indicates that most of them are also Nb carbonitrides (Fig. 3). The selected area diffraction (SAD) pattern shows that they have an fcc structure with lattice parameter a = 0.44 nm (Fig. 3d). The size range of such residual Nb(CN) particles is from 0.05 to 0.4 μm. They could precipitate during cooling after solidification and annealing and are referred to as secondary Nb(CN). Very occasionally, rod-like particles are found (Fig. 3e). EDS reveals that the particle is (Cr,Nb)-rich nitride (Fig. 3f). SAD (Fig. 3g) suggests that such particles have a tetragonal structure with lattice parameters of a = 0.30 nm and c = 0.74 nm. These results demonstrate that the (Cr,Nb)-rich nitride is Z-phase (CrNbN) [7]. No other type of precipitate was found. It should be noted here that a few dislocations were observed, which could have been induced by water quenching. 3.2. Microstructural evolution during aging The aging characteristics of Alloy 709 were investigated over the temperature range 550–750 °C. The subsequent sections report the influence of time and temperature on the phase transformations. Each observed phase was characterised in detail. 3.2.1. Aging at 550 °C In the samples aged at 550 °C there was no evidence for any other particles beyond what was found in the as-received material, either at grain boundaries or in the grain interiors (Fig. 4). This finding suggests that the mechanical behaviour of the Alloy 709 is not expected to degrade too much during service at 550 °C. However, aging did promote segregation of Cr and Mo at grain boundaries (Fig. 4). The Cr and Mo segregation to GBs has been widely reported in steels [8–11]. 3.2.2. Aging at 650 °C Aging at 650 °C significantly modified the microstructure of the Alloy 709. 500 h aging resulted in the formation of globular particles along the grain boundaries and of plate shaped particles in the grain interiors and at incoherent twin boundaries (ITB) (Fig. 5). EDS maps indicate that the grain boundary particles contained predominantly Cr,

3.2.3. Aging at 750 °C Aging at a higher temperature (750 °C) produced coarse particles at grain boundaries, as can be seen from Fig. 13. EDS mapping indicates that the GB particles consist mainly of M23C6 carbides and a Si-

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(b)

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Fig. 11. (a) BF-STEM image of particles at grain boundaries (GBs) in the sample aged at 650 °C for 2000 h, (b) corresponding EDS maps, showing that most GB particles are M23C6 carbides along with some Si-rich θ phases, (c) a superposition of Si and Cr-EDS maps shows that some θ phases grew from the globular M23C6 carbides (white arrow) and from the M23C6 plates (black arrow) (c). Note: the holes at the upper left side of the image are due to preferential etching of the particles during twin-jet polishing.

60 Fe Ni Cr Mo C

at %

40

20

0 0

20

40

60

80

100

120

nm

(b)

(a)

(c)

(d)

Fig. 12. (a) The plate-like particles formed on the CTBs in the sample aged at 650 °C for 2000 h, (b) EDS linescan across a particle revealing that the particle is (Cr,Mo)-rich carbide (i.e., M23C6), (c) confirmation by HRTEM and (d) very fine particles have nucleated on the M23C6 plate (inset is high magnification image of the arrowed plate of Z-phase)..

containing phase. Spot EDS spectra reveal that the Si-containing phase is a (Cr,Ni,Si,Mo)-rich nitride (Fig. 14b). SAD patterns taken from the Si-containing phase and matrix in grain 1 shown in Fig. 14a are illustrated in Figs. 14c-d, indicating that the phase has a cube-to-cube orientation relationship with the matrix, the same as that between an

M23C6 carbide and the matrix, with nearly three times the lattice parameter of the matrix. In the [001] diffraction pattern some of the diffraction spots (e.g. {200},{420}) are invisible. This is consistent with the forbidden reflection conditions h,k,l all even and h + k + l ≠ 4n, n integral for the diamond-cubic structure. This phase is thus identified as

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(a)

(b)

(c) Fig. 13. (a) BF and (b) HAADF STEM images of particles at GBs in the sample aged at 750 °C for 500 h, (c) corresponding EDS maps showing that most GB's particles are M23C6 carbides with some Si-rich θ phases.

the diamond-cubic structure θ phase with a = 1.07 nm, which is consistent with the Cr3Ni2SiN (θ phase) reported in the study of Sourmail and Bhadeshia [6]. Spot EDS analysis indicates that the θ phase contains some Mo and Fe (Table 3) and thus the general formula can be

expressed as (Cr,Mo)3(Ni,Fe)2SiN. It is interesting to note that some plates at CTBs show different contrast for different segments (Figs. 15a and b). EDS maps reveal that the left side (dark segment in Fig. 15b) of the plate is the θ phase

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2000

Fe Matrix θ

1600

Counts

Cr

Ni

1200

Fe Ni 800

Si

Cr 400

Mo Fe

N Cr

Ni

0 0

1

2

3

4

5

6

7

8

9

keV

(a)

(b)

(c)

(e)

(d)

Fig. 14. (a) A BF-TEM image of GB θ phase in the sample aged at 750 °C for 500 h, (b) spot EDS spectrum from matrix and the θ phase revealing that the θ phase is (Cr,Ni,Si,Mo)-rich nitride, (c) composite SAD patterns from grain 1 and θ phase were recorded along [114] (c), [101] (d) and [001] (e) revealing that the θ phase has a diamond-cubic structure and a cube-to-cube orientation relationship with the matrix. Table 3 Composition of (Cr,Mo)3(Ni,Fe)2SiN in Alloy 709 (at.%). Temperature, °C

Time, h

650 750

2000 500 1000 2000

Cr 35.0 31.7 35.0 34.7

± ± ± ±

Mo 1.8 1.1 3.2 2.3

Ni

8.8 ± 0.8 10.0 ± 0.3 10.7 ± 1.1 10.9 ± 0.8

whereas the middle part (bright segment in Fig. 15b) is M23C6 carbide and that Z phases nucleated on the M23C6 carbide (Fig. 15b). This observation suggests that the θ phase may nucleate on the end of the M23C6 plate. Further aging produced more θ phases at coherent twin boundaries, as shown in Fig. 16. In addition to the formation of platelike θ phases, some irregular θ phases were also observed, which

31.6 33.6 31.9 30.7

± ± ± ±

Fe 3.2 2.2 3.4 3.1

3.6 2.8 3.3 3.2

± ± ± ±

Si 1.8 0.6 1.0 0.6

19.6 20.9 18.0 19.7

± ± ± ±

3.1 2.5 1.3 1.4

nucleated on the shell (being M23C6 carbide) of the complex particles with the Nb(CN) core, as can be seen in Fig. 17. This probably indicates that the θ phase is more stable than the M23C6 carbide. It is worth noting that the core of some complex particles is still Nb(CN) (Fig. 17b), which means that, even after aging at 750 °C for 2000 h, the Nb(CN) is still stable and has not completely dissolved.

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(a)

(b)

(c) Fig. 15. (a) HAADF-STEM image from the sample aged at 750 °C for 1000 h shows plates at CTB, (b) high magnification image of the plate at CTB in Fig. 15a revealing different contrast for different segments and (c) EDS maps of the marked region in Fig. 15b indicate that the dark segment is the θ phase while the bright part is the M23C6 carbide and that Z phase had nucleated on the M23C6 plate.

EDS analysis (Figs. 18b and c) suggests that the particles on dislocations are Z-phase. Compared with the samples aged at 650 °C, aging at 750 °C promotes more Z-phase on dislocations (Fig. 18a). Very occasionally, it was found that a Z-phase particle ~10 nm in width nucleated on a M23C6 carbide (Figs. 18d and e). FFT of the HR-TEM image reveals that there is an orientation relationship between the Z-phase and matrix (Fig. 18f), i.e. (100)Z//(110)γ and [001]Z//[001]γ. The observations of the various phases are summarised in Table 4.

3.2.4. Cr depletion at grain boundaries Aging produced precipitation of Cr-rich M23C6 carbides at the grain boundaries, probably leading to Cr depletion in the vicinity of the grain boundaries and thus degrading the corrosion resistance of Alloy 709. EDS linescan was used to reveal Cr grain boundary depletion. Fig. 19 shows Cr profiles across grain boundaries in the samples aged at 650 °C. For example after aging for 500 h at 650 °C, an obvious depletion of Cr was found at the GB where the content of Cr is about 6.4 wt% (Fig. 19b1), which is much lower than the minimum of 10.5 wt% Cr for stainless steel. With increasing aging time, the degree of Cr depletion at

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Fig. 16. (a) HAADF-STEM image of the sample aged at 750 °C for 2000 h shows plates at twin boundaries, (b) EDS map of the marked region in Fig. 16a indicates the formation of θ phase at twin boundaries and in the twin interior and (c) a superposition of Si and Cr-EDS maps reveals that the plate shaped particle at the CTB is a complex particle containing M23C6 and θ phase.

GB alleviates but the width of the Cr-depleted zone (Fig. 19) increases. Even after 2000 h, the Cr depletion is not removed completely although the Cr content at the GB has recovered to ~12 wt% (Fig. 19b3). When the aging temperature increased to 750 °C (Fig. 20), although Cr depletion is still present in the sample aged for 500 h it is not as serious as that in the samples aged at 650 °C. This is because the calculated diffusivity of Cr at 750 °C is 30 times faster than that at 650 °C based on the data in ref. [12, 13]. Also, we can use an alternative method to describe the degree of Cr depletion, which is expressed by area fraction, %, = A1/(A1 + A2), where A1- integral above the depletion curve; A2- integral below the depletion curve, as shown schematically in Fig. 21. The calculated results based on the Cr profiles shown in Figs. 19 and 20, are given in Table 5, indicating that the degree of Cr depletion (%) decreases with increasing aging time and temperature, for instance, only 4.6% for the sample aged at 750 °C for 500 h.

frequently observed precipitate in this study, with a variety of shapes and locations. For instance, globular M23C6 was observed preferentially at GBs at 650 and 750 °C while plates grew initially at ITBs. Further blocky and plate shaped M23C6 particles formed in the grain interiors. The blocky M23C6 particles in the grain interior nucleated on Nb(CN) (Figs. 6 and 7) or dislocations (Fig. 8). With increasing aging time and temperature, plates of M23C6 start to form at CTBs and part of the GB M23C6 particles are replaced by the θ phase, although M23C6 remains the major phase at GBs, even after long-term aging. The Cr and Mo content of the M23C6 carbides increases with increasing aging time or temperature (Table 2), which suggests that Cr and Mo stabilise in M23C6. M23C6 is commonly observed in high Cr steels [18–20] even after short term aging. Indeed, in our tensile samples (subjected to ~ 30 mins holding time) for as-received material, some fine M23C6 carbides were observed at GBs at 650 °C while a nearly continuous distribution of M23C6 particles at GBs was found at a testing temperature of 750 °C [21]. This indicates that the growth rate of M23C6 phase is much higher at 750 °C. A previous study shows that M23C6 precipitates preferentially at GBs in short times during service at 600–1000 °C [22]. Sourmail observed the formation of M23C6 after very short aging times (e.g. 30 min) at 750 °C in a stabilised stainless steel [7]. The TTP curve of M23C6 in TP310 stainless steel exhibits a C-shape with a nose temperature of ~750 °C [23]. These findings are in accordance with our observations.

4. Discussion 4.1. Behaviour of various precipitates 4.1.1. Nb(CN) In this study the initial microstructure of the Alloy 709 prior to aging contains residual Nb(CN). This is consistent with previous studies on 20/25/Nb-stabilised steels which noted that Nb(CN) precipitated preferentially [14–17]. Several previous studies [14,15] demonstrated that the first phase to form during aging was Nb(CN) and this phase is regarded as the major short term precipitate. As the distribution of Nb (CN) particles is not uniform (as seen in Figs. 3 and 4), its average number density from different regions is used to reveal if aging leads to the precipitation of Nb(CN). The measurement indicates that aging indeed promotes the formation of Nb(CN) (2.6 ± 0.4 × 1011 m−2 for the as-received sample vs 4.1 ± 0.5 × 1011 m−2 for the one aged at 550 °C for 500 h). EDS mapping of the samples aged at 650 or 750 °C also showed that the fine Nb-rich phases (assumed to be Nb(CN)) were easily observed at the grain boundaries but rarely in the as-received sample. Both findings suggest that aging produced precipitation of Nb (CN), which is linked to the condition of the as-received material. As the as-received sample had received a solution treatment at 1100 °C, the subsequent quench produced C, N and Nb supersaturation. Consequently, on aging (e.g. at 550–750 °C), these elements are rejected from solution, to form Nb(CN). Additionally, the Nb(CN) particles act as nucleation sites for M23C6 carbides (Figs. 6 and 7).

4.1.3. θ (Cr3Ni2SiN) phase θ (Cr3Ni2SiN) phase was observed to occur at similar locations to the M23C6 phase, and to have a similar morphology, thus making it hard to distinguish one from the other without composition or careful diffraction analysis. Cr3Ni2SiN phase was found in Alloy 709 aged at 650 °C for 1000 and 2000 h and at 750 °C. Its actual composition includes substantial amounts of Mo and Fe, and the Mo concentration increases with temperature (Table 3). Thus a more general formula (Cr,Mo)3(Ni,Fe)2SiN is proposed. Very few published studies report its presence. Sourmail and Bhadeshia [6] demonstrated that (Cr3Ni2SiN) phase with some Mo and Fe was generated after aging NF709 and a NF709 variant at 750 and 800 °C. They proposed that θ (Cr3Ni2SiN) phase is a particular composition of the η structure whose lattice parameter is 1.062 nm (JCPDS 17-330), which means it has an extremely similar lattice parameter to that of M23C6, but the latter is cubic while the former is diamond-cubic. Early work also reports this formula. For instance, Horne reported the occurrence of a Cr3Ni2SiN in 21Cr40Ni0.22 N0.02C steel [24]. A nitrogen rich similar phase (Cr3Ni2SiN) is found after furnace aging of a 20Cr25Ni5Mo0.2 N steel [25]. This phase is reported after both 5 and 300 h at 850 °C and therefore is probably an equilibrium phase. This point is also supported by our observations: the phase was found at longer aging and higher aging temperature and its Mo content increases with temperature

4.1.2. M23C6 phase In general, aging of alloy compositions with stabilisation ratios [wt %Nb/wt%(C + N)] below 7.7 (stoichiometric composition) should result in the precipitation of M23C6. In Alloy 709 the Nb/(C + N) ratio is 1.1 and thus M23C6 is generally expected. Indeed, M23C6 is the most

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(a)

(b)

(c) Fig. 17. (a) A HAADF-STEM image from the sample aged at 750 °C for 2000 h shows some composite particles, (b) EDS maps of the region marked in Fig. 17a and (c) a superposition of Nb, Cr and Si-EDS maps revealing that some θ phases grew on the composite particles. 418

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(a)

(b)

1200

Cr

Fe

Fe

matrix Z phase

1000

Ni

Counts

800 600

Cr

400

Nb Mo

N

200

Cr

Si 0 0

1

2

3

4

5

6

keV

(c)

(d)

(e)

(f)

Fig. 18. (a) A BF-STEM image taken from the sample aged at 750 °C for 500 h shows many Z phase particles on the dislocations, (b) EDS maps and (c) spot EDS spectra (c) confirmed that they are (Cr,Nb)-rich nitrides (Z phase), (d) BF-TEM image from the sample aged at 750 °C for 2000 h shows the Z phase on the M23C6 carbide, (e) HRTEM image of the Z-phase and (f) FFT of the HRTEM image reveal an OR between the Z-phase and the matrix of (100)Z//(110)γ and [001]Z //[001]γ.

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Table 4 Summary of observations of secondary phases. Phase

Location/morphology

Typical size

Nb(CN)

Anywhere (rectangular) Anywhere (globular) Grain boundary (globular) CTB and ITB (plates) Around Nb(CN) (cuboidal) Grain boundary (globular) CTB and ITB (plates) Around M23C6 (cuboidal) Dislocation (rods) On M23C6 (rod)

1 to 10 μm 50 to 400 nm 100 to 500 nm

M23C6 (Cr,Mo)3(Ni,Fe)2SiN Z-phase

(Table 3). Those steels where (Cr3Ni2SiN) phase was observed have a high nitrogen content (Table 6), which indicates that this phase is stabilised by nitrogen. It should be mentioned here that G phase (Ni16Nb6Si7) is an alternative Si rich phase which has been observed in some 20Cr25Ni stabilised steels. For example, in 20Cr25Ni Nb and C containing steels, Powell et al. [4] and Ecob et al. [3] observed G phase and found that NbC transforms partially to G phase over time. However, the G phase was not found in this study, NF709 and a NF709 variant [6]. We also did not find evidence for Nb(CN) partially transforming to the other phase. This discrepancy could be associated with composition. For instance, compared with Alloy 709 and NF709, the steel in the study by Powell et al. has high Nb and Si but low C (Table 6), with a stabilisation ratio [wt%Nb/wt%(C + N)] = 14.5, [4] much larger than 7.7 (stoichiometric composition). In other words, after the first Nb(CN) precipitates, the steel still has surplus Nb but lacks C. The surplus Nb is the prerequisite for the formation of Nb-rich phase (e.g. G phase). Additionally, carbon could influence the formation of G phase. Indeed, Titchmarsh and Williams noted [26] and provided evidence [27] that the G phase formed preferentially only when carbon was not available. Thus, precipitation behaviour in 20Cr25Ni-based steels is linked with their minor constituents. For instance, (Cr,Mo)3(Ni,Fe)2SiN phase is found in high N-containing steels but G (Ni16Nb6Si7) phase is found in steel with high Nb and low N and C (Table 6). Now the question arises as to whether M23C6 or θ (Cr3Ni2SiN) phase is the more stable. No study has dealt with which phase is more stable. However, this study gives us some hints. The former phase was observed in the earlier stages, which indicates that θ (Cr3Ni2SiN) phase is probably the more stable. On the other hand, in addition to the θ phase formed on M23C6 phase (Figs. 11c and 17), complex rods, where the middle segment is M23C6 while the side segment is θ phase, were observed (Fig. 16), which could result from two possibilities: the θ phase grew on M23C6, or M23C6 transformed partially to θ phase. The EDS mapping shown in Fig. 16 can show whether the transition from M23C6 to θ phase occurs. As θ is enriched in Si while M23C6 contains higher Cr (Tables 2 and 3), M23C6 and θ phases can be distinguished by the distribution of these two elements. Fig. 16 shows clearly that the rod (arrowed in Fig. 16a) at the CTB is mixed M23C6 and θ phases; the M23C6 phase in that rod does not appear as a complete rod shape as observed in other regions (Fig. 16); some M23C6 phases were wrapped by θ phase. These findings indicate that a transformation of M23C6 to θ

100 to 500 nm 5 to 30 nm

phase must occur. We can thus conclude that θ phase is more stable than M23C6 in this study. 4.1.4. Z-phase Z-phase is a nitride with the formula CrNbN which usually forms in Nb-stabilised austenitic stainless steels with a high content of nitrogen [7]. Very occasionally rod-like Z-phase was found in the as-received material (Fig. 3) where it is present as residual particles, that is coarse particles formed during manufacture, as observed in NF709 [6]. It has been reported that Z phase can form at high temperature: Raghavan et al. [28] found its formation during annealing (1 h at 1027 °C) of an 18–12 steel containing 0.3wt%Nb and 0.09wt%N. The work of Robinson and Jack [29] indicates that the solvus of the Z phase is between 1300 and 1350 °C in a steel containing large amounts of Nb and N. In this study, Z phase does not form at 550 °C (Fig. 4) but starts to precipitate during aging at 650 °C (Figs. 8d and 10d). The precipitates are fine (~ 10 nm), they are nucleated on dislocations and their number increases with aging time and temperature. These observations are in a good agreement with the literature. For instance, Robinson and Jack [29] observed the formation of Z phase between 700 and 1000 °C in a 20Cre9Ni steel containing 0.38 wt%N and 0.27 wt%Nb. Vodárek [30] investigated the stability of Z phase in a type 316LN alloy: they had a mean size of 6 nm after 82 h at 650 °C and 12 nm after 37,890 h at the same temperature. The creep deformation behaviour of materials is generally dependent on the extent of matrix precipitation. Thus, it is expected that the precipitation of such stable fine Z-phase on dislocations could have a significant influence on the subsequent creep behaviour. 4.2. Effect on mechanical properties Aging at 550 °C leads to the segregation of Cr and Mo to grain boundaries whereas depletion of Cr is observed in the samples aged at 650 °C. Such a variation in the concentration of the elements influences not only the corrosion resistance of Alloy 709 but its mechanical behaviour because Cr and Mo also play a role in solid solution strengthening. The formation of nearly continuous GB particles, especially of the predominant phase - Cr-rich M23C6 carbide, could lead to a decrease in ductility of aged Alloy 709. The details of effect of aging on tensile

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Fe Ni Cr Mo

60

wt %

40

20

0 0

500

1000

(a1)

1500 nm

2000

2500

3000

(b1) Fe Ni Cr Mo

60

wt %

40

20

0 0

500

1000

1500 nm

2000

2500

3000

(b2)

(a2)

Fe Ni Cr Mo

60

wt %

40

20

0 0

500

1000

1500 nm

2000

2500

3000

(b3)

(a3)

Fig. 19. (a) HAADF-STEM images and (b) EDS linescans along the lines shown in the images taken from samples aged at 650 °C for 500 h (1), 1000 h (2) and 2000 h (3). Depletion of Cr and Mo near the grain boundaries was observed. The grain boundary Cr content gradually increases with increasing aging time. Note: two troughs (arrowed) in Fig. 19(b1) are because the linescan crosses two M23C6 particles; the dashed line in Fig. 19 (3b) represents the Cr content in the matrix without Cr depletion.

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Fe Ni Cr Mo

60

wt %

40

20

0 0

500

1000

(a)

1500 nm

2000

2500

(b)

Fig. 20. (a) HAADF-STEM image taken from the sample aged at 750 °C for 500 h and (b) EDS linescan along the line in Fig. 20a, showing that the depletion of Cr and Mo has nearly disappeared. Note: two small peaks (arrowed) in Fig. 20b are due to the linescan passing across two M23C6 particles. The dashed line in Fig. 20b represents the Cr content in the matrix without the Cr depletion.

Cr, wt.%

Table 5 The degree of Cr depletion calculated for the profiles shown in Figs. 19 and 20.

A1

Aging condition

500 h at 650 °C

%

27.2

1000 h at 650 °C 24.9

2000 h at 650 °C

500 h at 750 °C

19

4.6

Note: % = A1/(A1 + A2), A1- integral above the depletion curve; A2- integral below the depletion curve, as shown schematically in Fig. 21.

Cr profile across GB

the as-received state and after static aging at 550, 650 and 750 °C has been investigated in detail using TEM. The following conclusions could be drawn from this study:

A2 Distance

1. The prominent precipitate in as-received Alloy 709 is Nb(CN). Very occasionally, rod-shaped Z phase (CrNbN) was found. 2. No significant change in microstructure was observed at 550 °C after aging even up to 2000 h but the segregation of Cr and Mo to grain boundaries was found. No microstructural evolution means that Alloy 709 is fairly stable at 550 °C. 3. Aging at 650 °C produced a nearly continuous distribution of globular M23C6 phases at grain boundaries, plate-like M23C6 phases at twin boundaries and in the grain interior, and blocky M23C6 phases on Nb(CN). Fine dispersoid Z phases on dislocations were observed after aging for 500 h; their amount increases with aging time. 4. θ ((Cr,Mo)3(Ni,Fe)2SiN) phase was generated at grain boundaries after aging for 1000 h at 650 °C. After aging at 750 °C, θ phase nucleates on M23C6 carbides and evidence for the transformation of M23C6 to θ phase was found, which suggests that θ phase is the more stable at this temperature.

Fig. 21. Shows schematically the region of Cr depletion. Note: A1- integral above the depletion curve, A2- integral below the depletion curve.

behaviour were given in ref. [21]. For instance, aging at 650 °C for 2000 h resulted in a one-third reduction in elongation at room temperature compared with the as-received coupons. M23C6 carbide is a brittle phase [31] and cannot accommodate deformation, thus causing M23C6 carbide cracking during deformation (Fig. 22a) and leading to intergranular failure (Fig. 22b) [21]. However, aging at 550°C did not promote the formation of new phases and elemental (e.g. Cr) depletion at grain boundaries and thus it is expected that Alloy 709 will not be degraded significantly during service at 550°C. 5. Conclusions The microstructure (especially the secondary phases) of Alloy 709 in

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Table 6 Composition (wt%) of the steels investigated in some of the studies quoted. In study Alloy 709 NF709 Jargelius-Pettersson [25] Powell et al.

Cr

Ni

Mn

Mo

Si

Nb

Ti

N

C

19.69 20.28 19.8 19.4

25.0 24.95 25.0 24.4

0.88 1.00 1.44 0.74

1.46 1.50 4.59 –

0.28 0.41 0.54 0.61

0.23 0.26 – 0.68

< 0.01 0.05 – –

0.14 0.167 0.21 0.01

0.063 0.06 0.014 0.037

(a)

(b)

Fig. 22. (a) SEM image of the longitudinal cross-section of room temperature tensile coupons subjected to aging at 650 °C for 2000 h, revealing M23C6 carbide cracking (arrowed), (b) SEM image of fracture surface of room temperature tensile coupons subjected to aging at 650 °C for 2000 h, showing intergranular failure [21].

Acknowledgements

[8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21]

This study is part of a project funded by the Research Council of United Kingdom (RCUK) award EP/N016351/1 and United States of America's Department of Energy (DOE) nuclear energy university program (NEUP) award number 2015-1877/DE-NE0008451. We thank Prof. Ian Jones for useful discussions. Data availability The data that support the findings of this study is available from the corresponding authors on request.

[22] [23] [24] [25] [26]

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