phase separation study of in-service thermally...

17th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors

August 9-13, 2015, Ottawa, Ontario, Canada

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PHASE SEPARATION STUDY OF IN-SERVICE THERMALLY AGED CAST

STAINLESS STEEL – ATOM PROBE TOMOGRAPHY

Martin Bjurman1, Mattias Thuvander2, Fang Liu2, Pål Efsing3.

1Studsvik Nuclear AB / Royal Institute of Technology (KTH), Nyköping, SE-611 82, Sweden

2 Chalmers University of Technology, Department of Applied Physics, Goteborg, SE-412 96, Sweden

3Ringhals AB / Royal Institute of Technology (KTH), Väröbacka, SE-430 22, Sweden

ABSTRACT

Embrittlement of Duplex Stainless Steels by thermal aging shortens the service life of structural

components in LWRs. This is an important issue when life extension programs are aiming at 60-

80 years in service. Cast and welded austenitic stainless steels, which contain some ferrite, are

known to be affected by thermal aging. Historically, many LWR components of complex

geometry have been cast in the Mo-containing quality CF8M. Aging is attributed to two types of

phase transformations; Demixing of the ferrite by spinodal decomposition into Cr-rich ´ and

Fe-rich regions; and precipitation of G-phase, carbides and other secondary phases.

A study was conducted on two in-service aged large casting CF8M elbows exposed for 72 kh at

291ºC and 325ºC, respectively, followed by 22 kh at a reduced service temperature. Atom Probe

Tomography was used to characterize the decomposition of the ferrite for both aging states.

Spinodal decomposition and nucleation of precipitates, i.e. G-phase, have been identified. The

extent of phase transformation increases with exposure temperature, and the mechanical

properties follow the same trend.

Keywords: Thermal Aging, Atom Probe Tomography (APT), Cast Stainless Steel (CASS), CF8M

1. INTRODUCTION

Thermal aging embrittlement of cast and welded “Austenitic“ Stainless Steels (CASS) for nuclear power

plant applications has been extensively studied during the 80´s and early 90´s. The thermal aging leads to

an increase in hardness and tensile strength, and a decrease in ductility, impact strength, and fracture

toughness. A renewed interest has arisen in recent years for these issues driven by life extension programs

and a difficulty in predicting the behavior and microstructure after 60-80 years at reactor temperatures.

Typical compositions of cast stainless steels used in Nuclear Power Plants (NPPs) are presented in Table

1.

These materials have a duplex solidification microstructure consisting of austenite and δ−ferrite phases.

The main aging phenomenon is the thermal diffusion driven decomposition of the ferrite into iron-rich -

phase and chromium-rich ´-phase due to the miscibility gap in the Fe-Cr phase diagram [1-3]. A

maximum rate of decomposition occurs at 475ºC, hence the name 475ºC-embrittlement. The thermal

aging phenomenon of ferritic grades from spinodal decomposition in high temperature applications of

280-500°C is widely known [4-9]. Research has focused on investigating both binary alloys and

commercial grades used in NPPs. An additional contribution to embrittlement comes from the

precipitation of G-phase, rich in Ni, Si, Mn, Mo and Ti (ideally Ni16Ti6Si7), in the δ−ferrite and

precipitation and preferential growth of carbides and nitrides at the ferrite/austenite interface. These

effects change the material's mechanical properties and to some extent the corrosion behavior [10],



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though consensus has not been reached on the magnitude and interaction with mechanical properties.

Studies on cast alloys have revealed that the phase boundary carbides play a significant role in thermal

embrittlement at temperatures greater than 400°C, but have insignificant effect on the embrittlement at

exposure temperatures below 400°C [9]. Also, aging at 400°C results in spinodal decomposition of the

ferrite phase which strengthens the ferrite phase and increases e.g. cyclic hardening. Thermal aging at

465°C results in the nucleation and growth of large α’ particles and other phases such as sigma phase,

which do not change the tensile or cyclic hardening properties of the material.

2. MATERIAL

Test materials are extracted from the Nuclear Power Plant Ringhals 2 steam generator (SG) loop 2 inlet

(hot) and crossover (outlet) to reactor coolant pump (cold) elbows. The material is ASTM 351 CF8M

with chemical composition according to Table 2. A slightly higher Cr- and Cu-content is seen for the hot-

leg. The tabulated ferrite contents were calculated from the composition using Schaeffler diagrams. Heat

treatment of the cast elbows was originally conducted at 1050ºC for 24 h followed by rapid cooling in

accordance with the standard. The microstructure consists of regions with both columnar and equiaxed

solidification structures. Local δ−ferrite contents measured by ferritescope vary from 1.5% to 22.5% [11].

Ferrite contents from the regions where specimen were extracted exhibited close to 10% ferrite and

equiaxed structure. Aging times and temperatures are presented in Table 3 with two exposure periods at

different temperatures due to power reduction during the last 22kh of service.

The ferrite structure seen in figure 1 exhibits large ferrite areas from the slow cooling of these large castings.

The microstructure in the chosen regions is equiaxed, with ferrite spacing of approximately 100 m and

some carbide precipitates are present in the cross-section. No large microstructural differences are seen

between the castings.

3. EQUIPMENT

3.1 Scanning Electron Microscopy

A tungsten filament SEM, JEOL 6300 with a Thermo-Fischer NS-7 EDS was used for elemental analysis

of ferrite and austenite phases respectively. EDS-analyses were carried out at 20 keV and averaged over

several grains.

3.2 Focused Ion Beam Scanning Electron Microscopy

Samples for APT were prepared using FIB/SEM lift-out technique [12] using a FEI Versa 3D DualBeam.

It was necessary to do site-specific preparation as the ferrite volume fraction is rather low. An image with

the selected region for APT analysis of hot-leg is shown in figure 2.

3.3 Atom Probe Tomography

The instrument used is a local electrode atom probe, Imago LEAP 3000X HR. Analyses were carried out

both in voltage pulse mode and in laser pulse mode, with very similar results. For voltage pulsing, the

sample temperature was 70 K and the pulse fraction was 15% of the DC voltage, whereas the laser

pulsing was carried out at 50 K with a pulse energy of 0.30 nJ (wavelength 532 nm). The pulse frequency

was 200 kHz in all analyzes. The instrument has a detection efficiency of 37%.

4. RESULTS

4.1 EDS-analysis

The results of EDS-analysis of ferrite and austenite respectively are presented in Table 4. Values fit well

with results from charge data, note tabulated at.% of EDS and Wt% of charge data. Further, the higher Cr



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and lower Ni content of hot-leg is seen in the results for both ferrite and austenite. The ferrites'-Ni content

is fairly low compared to literature data.

4.2 Atom Probe Tomography

The composition of the ferrite in the hot-leg and the crossover-leg, respectively, is presented in Table 5.

The differences between the two materials are small, and must be regarded as insignificant, perhaps with

the exception of Mo.

The APT analyses of the hot-leg showed clearly the phase separation into α and α' phase, as well as the

formation of G-phase precipitates. In the crossover-leg, the microstructural changes are much weaker.

The two materials are compared using the radial distribution function (RDF), see curves in figure 3 and 4.

The RDFs represent radial concentration profiles evaluated starting from each detected atom of the

specific element. Plots are given as the bulk normalized probability density of finding an atom of the

chosen type at radial distance r from each equal atom respectively. The composition of the three phases

was determined for the hot-leg using iso-concentration surfaces with thresholds listed in Table 6. Hence,

each volume fulfilling certain compositional criteria is identified and iso-concentration surfaces enclosing

these volumes are created. An iso-concentration plot of the hot-leg is presented in figure 5, where a box

(37×34×34 nm3) within the analyzed volume is plotted with red α-, blue α'-, and green G-phase. Atomic

maps of projected volumes (20×20×5 nm3 slices) for the hot-leg and crossover-leg are presented in

figures 6 and 7, respectively. From Cr-concentrations maps, it is seen that spinodal decomposition occurs

in both hot- and crossover-leg, but substantially less in the crossover leg.

The wavelength of the spinodal decomposition of the hot-leg was determined from the Cr RDF curve to

6.8 ± 0.2 nm. The wavelength of crossover-leg is due to the limited decomposition more difficult to

quantify, but is estimated to be 4 ± 1 nm. Also the amplitude of the spinodal decomposition was

determined to 10.7 at.% for the hot-leg and 8.1 at.% for the crossover-leg, respectively, following the

approach in [13].

The number density of G-phase precipitates in the hot-leg was determined to be (3.9 ± 0.5) × 1024 m-3.

From this value the characteristic distance between precipitates can be estimated as one over the cubic

root giving a length of 6.4 nm, which is close to the wavelength of 6.8 nm of the spinodal decomposition.

The average diameter of the G-phase precipitates was estimated to about 3.2 nm, and the volume fraction

was about 6%. It should be noted that the measurement of average size and volume fraction using APT is

less accurate than the measurement of number density, as there is a large influence of the threshold used.

In this case the threshold used was Ni+Si+Mn >15 at.%.

A proximity-histogram of the G-phase precipitates of the hot-leg is shown in figure 8, obtained using a

threshold of Ni+Mn+Si>20%. Included in the proximity-histograms are also Cu and P, showing increased

concentrations of these species in G-phase.

From the crossover-leg Ni-Si- and Ni-Mn-RDFs, presented in figures 9 and 10, there is a clear positive

interaction. In order to estimate the number density, size and composition of clusters in the crossover-leg,

an established cluster algorithm was applied [14]. The clustering of Ni, Si and Mn was evaluated using

the parameters dmax =denvelope=derosion=0.40 nm and Nmin=10. The number density of clusters in the

experimental dataset was 1.5×1025 m-3, compared to the number density of 9.5×1024 m-3 in a randomized

dataset, giving a difference of 5.6×1024 m-3. The average cluster contained only 21 atoms, with a

maximum of 102 atoms. The average composition of the clusters is presented in Table 7.

5. DISCUSSION

Chemical composition and ferrite morphology strongly affect the extent and kinetics of embrittlement

[15, 16]. Microstructure analyses of different cast and welded structures [17] show that the amount of



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ferrite in cast SS mainly depends on the composition, but also on the cooling rate. The size, distribution,

and morphology of the ferrite within the austenite matrix also depend on the solidification conditions

during the casting process.

Using the equations of Suutala [18] and the chemical composition in Table 2, the solidification route

appears to lie quite close to the limit between two-phase solidification,

LiquidLiquidLiquidand pure solidification compared to typical large CF8M-

castings. Hence the ferrite is either formed during solidification or transformed from austenite. Most

castings exhibit pure solidification. Calculations from the compositions of selected castings in Table 8

show that the actual solidification routes vary. The limit between the mentioned solidification routes is

located at Creq/Nieq of 1.95. This is highly sensitive to the C- and N-concentrations, so even compositional

changes within the measuring uncertainty would be sufficient to change the solidification route.

These castings are then annealed at 1050 ºC after casting followed by quenching. Quenching rates have

previously been shown to affect the rate of spinodal decomposition [19], which is partially why samples

are taken from approximately 10 mm below the surface of each leg. An indication of having significant

rate variations is the large local ferrite content variations of 1.5 - 22.5 %. A possible contributor to the

decomposition, especially of importance for the small decomposition of the crossover-leg, is

decomposition that might have occurred during cooling of the original casting [13].

Additional modelling using the Scheil equation and Thermo-Calc Calphad [20] modelling tools was made

to verify the solidification route. The results concur with the Suutala equations and show that these

approximate equations are fairly accurate in interpreting the solidification route. Both these techniques are

time independent, assuming diffusion of species to be infinite in the liquid phase, equilibrium at the

interface and zero diffusion in the solid phase. This is important to be aware of when comparing slowly

cooled heat-treated large castings with smaller lab-castings or single/multi-pass weld materials. In fact,

welds of these materials generally solidified through the liquidroute. The compositional

analysis of the ferrite in the investigated castings revealed a low Ni-content [21], which probably

suppresses spinodal decomposition and possibly increase G-phase formation rates [22].

The spinodal decomposition of the hot-leg both regarding wavelength (6.8 ± 0.1 nm) and amplitude

(10.7 at.%) agrees well with earlier studies [e.g. 23]. Investigations of CF8M aged below 300ºC are

scarce as the low degree of decomposition is difficult to measure.

Calculation of the aging equivalence between cold- and hot-leg using a modified Arrhenius equation

(equation 1) with activation energy Q = 243 (±80) kJ/mol [24], using neff = 0.16, the wavelength of the

crossover-leg would be 4.3 nm. The values of the activation energy vary with temperature and values

between Q = 230 kJ/mol, which is expected from Cr-diffusion, to 260 (±50) kJ/mol [25] have been

reported for relevant conditions.

𝑡 2 = 𝑡1exp⌊𝑄

𝑅(1

𝑇1−

1

𝑇2)⌋ (1)

Where t1 and t2 are equivalent aging times, T1 and T2 respective temperatures and R is the ideal gas

constant. G-phase precipitation occurred on the -´ boundary as previously shown [e.g. 23] and the

evolution is considered to be highly linked to spinodal decomposition. This is driven by the

diffusion/rejection of species, e.g. Ni and Si, from the respective phase. Hence the close link between the

two processes [26].

The number density of G-phase in the hot-leg was (3.9 ± 0.5) × 1024 m-3 with an average diameter of 3.2

nm which in comparison indicates fairly large particles, hence the G-phase precipitation is already highly

developed. The crossover-leg just showed tendency towards G-phase precipitation, with a concentration

of 5.6×1024 m-3measured by clustering technique. The cluster size indicated is 21 detected atoms,



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corresponding to 57 actual atoms, which is below the number of 116 atoms needed for a complete unit

cell. The low levels of G-phase seen in the crossover-leg compared to the hot-leg, also in comparison to

the spinodal decomposition, would indicate a higher activation energy Q than for spinodal decomposition.

A lower Q of 140 ± 60 kJ/mol has previously been reported [25]. Compared to the spinodal

decomposition, G-phase indicates a smaller difference in concentrations, hence would indicate a high Q.

Using the correlations of [26] where G-phase precipitate radius and spinodal length were correlated to

evolve with the same effective time exponent neff (equation 2), would still cause particle diameters rG

larger than 1.5 nm.

𝑛𝑒𝑓𝑓 =𝑑ln(𝑟𝐺)

𝑑ln(𝑡) (2)

.

Measured approximate values from clustering indicate a size of 0.7 nm. This correlation between G-phase

and spinodal decomposition could not be seen here, where the crossover-leg shows even smaller G-phase

precipitates compared to the spinodal length. A plausible explanation is that the G-phase precipitation

only occurs after the spinodal has developed to a certain extent.

The composition of the G-phase in the hot-leg, Table 6, shows, as expected, enrichment in Si, Mn, Ni and

P and tendencies of an enrichment in Cr, W and Co. No increase of Mo or V is measured, but instead a

significant increase in the Cu-content. Copper has been seen in possible precursors of G-phase of super

duplex SAF2507 [27]. The composition of the G-phase clusters in the crossover-leg, Table 7, also shows

reasonable proportions of G-phase species.

Results of tensile, hardness and fracture mechanical data [12] showed significant reduction of mechanical

properties with aging temperature and compared to the “as-received” values. A significant reduction of

Charpy U impact values was observed at room temperature, see Table 9. The deterioration of the

mechanical properties of the crossover-leg is substantial while the APT-results still only show limited

spinodal decomposition and only traces of G-phase precipitation.

CONCLUSIONS

In-service thermal aging of ferrite from CF8M was investigated by APT for spinodal decomposition and

precipitation. Two specimens were analysed from hot- and crossover-legs in both laser and voltage pulse

modes. The spinodal decomposition of the hot-leg aged at 325°C was significant. From analysis of the

RDF curve a spinodal length of 6.4 nm and a Cr-concentration amplitude of 10.8 at % was deduced. The

crossover-leg aged at 291°C showed only limited spinodal decomposition. The amount of G-phase was

estimated to 3.9×1024 m-3 and diameter 3.2 nm for the hot-leg, whereas the crossover-leg only showed a

weak tendency towards G-phase precipitation with a cluster density of 5.6×1024 m-3 of approximately 57

atoms in each position and, much less than expected from the hot-leg content.

The reduction in fracture resistance previously measured in the crossover-leg is high compared to the weak

spinodal decomposition measured by APT.

ACKNOWLEDGEMENTS

The authors wish to express their sincere gratitude to Anders Jenssen (Studsvik), Thomas Barkar (KTH)

the reference group consisting of Peter Ekström (Swedish Radiation Safety Authority), Massimo Cocco

(Forsmarks Kraftgrupp), Mattias Coudret (Oskarshamns Kraftgrupp) and Björn Forssgren (Ringhals) for

fruitful discussions.

The work was funded by the Swedish utilities, Swedish Radiation Safety Authority and Studsvik.



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specific specimen preparation for atom probe tomography”, Ultramicroscopy 107 (2007) 131-

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Decomposition in Fe-Cr by Atom Probe Tomography and Radial Distribution Function

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TABLES

Table 1. Typical composition of standard cast stainless steels in nuclear applications.

Grade C Mn P S Ni Cr Mo

CF3 0.03 0.60 0.003 0.002 9 18 <0.5

CF8 0.057 0.62 0.003 0.002 8.5 20 2.21

CF8M 0.074 1.21 0.032 1.24 9.59 18.67 2.73

Table 2. Chemical composition and Schaeffler ferrite content of Ringhals 2 SG-elbows.

C Si Mn P S Cr Ni Mo Ti Cu Co N Ferrite

content

hot .037 1.03 .77 .022 .008 20.0 10.6 2.09 .004 .17 .040 .044 20.1

cross

over

.039 1.11 .82 .020 .012 19.6 10.5 2.08 .004 .08 .035 .037 19.8



8

Table 3. Full power temperature exposure of Ringhals 2 SG-elbows.

Full power time ~70 000h ~22 000h

Hot-leg 325ºC 303ºC

Crossover-leg 291ºC 274ºC

T 34ºC 29ºC

Total Full Power time 92 000h

Table 4. Results from EDS-analysis of ferrite and austenite in as received material.

Atom % Fe Cr Ni Mn Si Mo

Hot-leg Ferrite 63.1 26.4 5.57 0.82 2.00 2.05

Austenite 65.7 20.8 9.41 0.96 1.82 1.30

Crossover-leg Ferrite 62.9 26.5 5.69 0.87 1.91 2.11

Austenite 65.5 20.5 10.26 0.80 1.82 1.14

Table 5. Composition (at.%) of the ferrite phase measured by APT, balance Fe.

Si Mn Cr Ni P Mo W V Cu Co

Hot 2.22 0.63 24.4 5.71 0.06 1.84 0.02 0.06 0.02 0.05

Crossover 2.10 0.62 23.9 5.80 0.05 2.20 0.001 0.06 0.03 0.03

Table 6. Composition (at.%) of the different phases of the hot-leg, balance Fe.

Limit Si Mn Cr Ni P Mo W V Cu Co

α Fe>72 1.10 0.18 12.9 3.66 0.02 1.10 0.01 0.03 0.02 0.05

α' Cr>30 1.98 0.57 42.0 4.09 0.04 2.13 0.01 0.11 0.01 0.04

G Ni+Mn+Si>20 11.0 4.28 16.5 20.5 0.21 2.16 0.03 0.05 0.18 0.06

Table 7. Average composition (at.%) of clusters in the crossover-leg.

Si Mn Cr Ni Mo Fe

Crossover 15.2 4.6 8.9 45.7 0.9 24.5

Table 8. Ferrite content, ferrite number and expected solidification route.

Origin of info. Component

Ferrite

content

FN -

Creq/Nieq

Ringhals 2 SG hot-leg 20.1 1.96 Liquid → Liquid + δ → δ + γ



9

Ringhals 2 SG Crossover-leg 19.8 1.97 Liquid → Liquid + δ → δ + γ

C.Pareige JNM 2010 Ingot A 2.33 Liquid → Liquid + δ → δ + γ

C.Pareige JNM 2010 Ingot B 1.91 Liquid + δ→ Liquid + δ + γ→ δ + γ

Table 9. Charpy Impact test results in Joule (KCU) at 20ºC of Hot and Crossover-legs from [11].

Unaged Aged tangetial Aged axial

Crossover 137 69 71

Hot 141 33 31

FIGURES

Figure 1. Polished and etched samples of crossover (left) and hot-leg (right), showing ferrite content and

scattered carbides precipitates.

Figure 2. SEM picture of sample needle extraction position of hot-leg.

1



10

Figure 3. Radial density functions of Cr of hot and crossover-legs respectively.

Figure 4. Radial density functions of Ni of hot and crossover-legs respectively.

0,98

1

1,02

1,04

1,06

1,08

1,1

1,12

0 1 2 3 4 5 6 7 8 9 10

Bu

lk N

orm

aliz

ed C

on

cen

trat

ion

Radial distance (nm)

Hot-leg

Crossover-leg

0,95

1

1,05

1,1

1,15

1,2

1,25

1,3

1,35

1,4

1,45

0 0,5 1 1,5 2 2,5 3 3,5 4 4,5 5

Bu

lk N

orm

aliz

ed C

on

cen

trat

ion


Hot-leg

Crossover-leg



11

Figure 5. Iso-concentration plot of hot-leg with red α-, blue α'-, and green G-phase. The size of

the box is 34×34×37 nm3.

Figure 6 Atomic maps of projected volumes (20x20x5 nm slices), left hot-leg and right

crossover-leg. Ni (blue), Mn (red), Si (green).

8

0

-8

8

0

-8

8 0 -8 8 0 -8



12

Figure 7 Atomic maps of projected volumes (20x20x5 nm slices) of Cr-concentration, left hot-

leg and right crossover-leg.

Figure 8. Proximity-histogram of hot-leg G-phase based on Ni+Mn+Si>20% iso-concentration

surface. Solid dots follow the left scale bar, crosses the right scale bar.

0

0,2

0,4

0,6

0,8

1

1,2

1,4

1,6

1,8

2

0

5

10

15

20

25

-10 -8 -6 -4 -2 0 2

Co

nce

ntr

atio

n (

at.%

)

Distance from interface (nm)

Ni

Si

Mn

Cu

PC

on

cen

trat

ion

(at%

)

8

0

-8

8

0

-8

8 0 -8 8 0 -8



13

Figure 9. Radial density functions of Si with respect to Ni as reference point.

Figure 10. Radial density functions of Mn with respect to Ni as reference point.

0,9

1

1,1

1,2

1,3

1,4

1,5

1,6

0 0,5 1 1,5 2 2,5 3 3,5 4 4,5 5

Bu

lk N

orm

aliz

ed C

on

cen

trat

ion


Hot-leg

Crossover-leg

0,9

1

1,1

1,2

1,3

1,4

1,5

1,6

1,7

1,8

1,9

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phase separation study of in-service thermally...

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