Informes Técnicos Ciemat 1048Diciembre, 2004

Departamento de Fusión Nuclear

Reduced ActivationFerritic/Martensitic SteelEurofer´97 as Possible StructuralMaterial for Fusion Devices,Metallurgical Characterization onAs-Received Condition and afterSimulated Service Conditions

P. FernándezA.M. LanchaJ. LapeñaM. SerranoM. Hernández-Mayoral

Reduced Antivation Ferritic/Martensitic Steel Eurofer´97 as Possible StucturalMaterial for Fusión Devices. Metallurgical Characterization on As-Received

Condition and after Simulated Service Conditions.

Fernández, P.; Lancha, A.M.; Lapeña, J.; Serrano, M.; Hernández-Mayoral, M.

73 pp. 45 figs. 62 refs.

Abstract:

Metallurgical Characterization of the reduced activation ferritic/martensitic steel Eurofer´97, on as-receivedcondition and after thermal ageing treatments in the temperature range from 400o to 600o for periods up to10.000 h, was carried out. The microstructure of the steel remained stable (tempered martensite withM23C6 and MX precipitates) after the thermal ageing treatments studied in this work. In general, thisstability was also observed in the mechanical properties. The Eurofer´97 steel exhibited similar values ofhardness, ultimate tensile stress, 0,2% proof stress, USE and T0, regardless of the investigated materialcondition. However, ageing at 600o C for 10.000 h caused a slight increase in the DBTT, of approximately23o. In terms of creep properties, the steel shows in general adequate creep rupture strength levels forshort rupture times. However, the results obtained up to now for long time creep rupture tests at 500o Csuggests a change in the deformation mechanism.

Acero Ferrítico/Martensítico de Activación Reducida Eurofer´97 Candidato a MaterialEstructural del Futuro Reactor de Fusión Nuclear. Caracterización Metalúrgica en

Estado de Recepción y Envejecido Térmicamente.

Fernández, P.; Lancha, A.M.; Lapeña, J.; Serrano, M.; Hernández-Mayoral, M.

73 pp. 45 figs. 62 refs.

Resumen:

En este informe se presentan los resultados de la caracterización metalúrgica del acero ferrítico/martensíticode activación reducida Eurofer´97. Los estudios se han realizado en estado de recepción y en el materialenvejecido térmicamente simulando las condiciones de operación del reactor, en el rango de temperaturasde 400oC - 600oC hasta periodos de 10.000 horas. Los resultados microestructurales han mostrado unaalta estabilidad estructural (martensita revenida con precipitados del tipo M23C6 y MX) en el rango detemperaturas investigado en este trabajo. En general, esta estabilidad también se ha observado en laspropiedades mecánicas. El acero Eurofer´97 presenta valores similares de dureza, resistencia a la trac-ción, límite elástico, USE y T0 independientemente del estado del material. Sin embargo, el envejecimientoa 6000C durante 10.000 horas ha causado un ligero incremento del DBTT en 23oC. Respecto a laspropiedades de fluencia, el Eurofer´97 presenta adecuados valores de resistencia a la fluencia para tiem-pos de ensayo cortos. Sin embargo, los resultados de los ensayos de larga duración obtenidos hasta elmomento parecen mostrar que a 5000C se está produciendo un cambio en el mecanismo de deformación.

CLASIFICACIÓN DOE Y DESCRIPTORES

S36

MARTENSITIC STEELS; METALLURGY; PHYSICAL METALLURGY; AGING;THERMOCHEMICAL TREATMENTS; HARDNESS; TENSILE PROPERTIES; MECHANICALPROPERTIES; TEMPERATURE RANGE 0400-1000K; THERMONUCLEAR REACTORS

REDUCED ACTIVATION FERRITIC/MARTENSITIC STEEL EUROFER'97

AS POSSIBLE STRUCTURAL MATERIAL FOR FUSION DEVICES.

METALLURGICALL CHARACTERIZATION ON AS-RECEIVED

CONDITION AND AFTER SIMULATED SERVICE CONDITIONS.

P. Fernández, A.M. Lancha, J. Lapeña, D. Gómez-Briceño,

M. Serrano, M. Hernández-Mayoral.

Reduced Activation Ferritic/Martensitic Steel Eurofer'97 as Possible Structural

Material for Fusion Devices. Metallurgical Characterization on As-Received

Condition and After Simulated Service Conditions.

P. Fernández, A.M. Lancha, J. Lapeña, M. Serrano, M. Hernández-Mayoral. Abstract Metallurgical characterization of the reduced activation ferritic/martensitic steel

Eurofer'97, on as-received condition and after thermal ageing treatments in the

temperature range from 400°C to 600° for periods up to 10000 h, was carried out. The

microstructure of the steel remained stable (tempered martensite with M23C6 and MX

precipitates) after the thermal ageing treatments studied in this work. In general, this

stability was also observed in the mechanical properties. The Eurofer'97 steel exhibited

similar values of hardness, ultimate tensile stress, 0.2% proof stress, USE and T0, regardless of the investigated material condition. However, ageing at 600°C for 10000 h

caused a slight increase in the DBTT, of approximately 23°C. In terms of creep

properties, the steel shows in general adequate creep rupture strength levels for short

rupture times. However, the results obtained up to now for long time creep rupture tests

at 500ºC suggests a change in the deformation mechanism.

Acero Ferrítico/Martensítico de Activación Reducida Eurofer’97 Candidato a

Material Estructural del Futuro Reactor de Fusión Nuclear. Caracterización

Metalúrgica en Estado de Recepción y Envejecido Térmicamente.

P. Fernández, A.M. Lancha, J. Lapeña, M. Serrano, M. Hernández-Mayoral.

Abstract

En este informe se presentan los resultados de la caracterización metalúrgica del acero

ferrítico/martensítico de activación reducida Eurofer’97. Los estudios se han realizado

en estado de recepción y en el material envejecido térmicamente simulando las

condiciones de operación del reactor, en el rango de temperaturas de 400°C-600°C hasta

periodos de 10000 horas. Los resultados microestructurales han mostrado una alta

estabilidad estructural (martensita revenida con precipitados del tipo M23C6 y MX) en

el rango de temperaturas investigado en este trabajo. En general, esta estabilidad

también se ha observado en las propiedades mecánicas. El acero Eurofer’97 presenta

valores similares de dureza, resistencia a la tracción, límite elástico, USE y T0

independientemente del estado del material. Sin embargo, el envejecimiento a 600°C

durante 10000 horas ha causado un ligero incremento del DBTT en 23°C. Respecto a

las propiedades de fluencia, el Eurofer’97 presenta adecuados valores de resistencia a la

fluencia para tiempos de ensayo cortos. Sin embargo, los resultados de los ensayos de

larga duración obtenidos hasta el momento parecen mostrar que a 500°C se está

produciendo un cambio en el mecanismo de deformación.

INDEX

1.- INTRODUCTION ............................................................................................................. 1

2.- EXPERIMENTAL PROCEDURE ....................................................................................

2.1 Material .......................................................................................................................

2.2 Microstructural Characterisation .................................................................................

2.3 Mechanical Properties .................................................................................................

3

3

3

4

3.- RESULTS ..........................................................................................................................

3.1 Chemical Composition ................................................................................................

3.2 Microstructural Characterisation .................................................................................

Optical and Scanning Electrón Microscopy ...............................................................

Phase Extraction and X-Ray Diffraction ....................................................................

Transmisión Electron Microscopy .............................................................................

3.3 Mechanical Properties .................................................................................................

Hardness Measurements .............................................................................................

Tensile Properties .......................................................................................................

Charpy Properties .......................................................................................................

Fracture Toughness ....................................................................................................

Low Cycle Properties .................................................................................................

Creep Properties .........................................................................................................

5

5

5

5

6

6

8

8

8

9

9

10

11

4.- DISCUSSION ................................................................................................................... 13

5.- SUMMARY AND CONCLUSIONS ................................................................................ 19

6.- REFERENCES .................................................................................................................. 20

7.- TABLES AND FIGURES ................................................................................................. 23

1

1. INTRODUCTION

Since the late 1970s Cr-Mo steels have been considered as potential first wall and breeding

blanket structural materials in fusion reactor systems. The fusion reactor applications require

steels resistant to radiation damage induced by bombardment from high energy neutrons

(14MeV), as well as steels that retains adequate toughness and elevated temperature strength

during service. The requirements for safe, routine operation, decommissioning of a fusion

plant and disposal of radioactive wastes have also demanded the development of steels with

enhanced radioactive decay characteristics. As a consequence of these requirements, new

steels denominated "reduced activation steels" were developed.

Basically reduced activation ferritic/martensitic steels were developed (1-7) by replacing Mo

in conventional Cr-Mo steels by W and/or V, and by replacing Nb by Ta. Alloy development

studies have shown that reduced-activation criteria of these steels can be achieved and

produced in the frame of the industrial viability. Special effort was made to lower the

restricted elements, emphasis was focused mainly on eliminating Nb because its concentration

in the steels must be

2

stability, favourable combinations of strength, toughness and resistance to radiation damage,

and superior characteristics, in terms of activation response, to their predecessors.

Subsequent efforts of the IEA collaborative international programme were focused on the

development and production of "large heats" of optimized steels with similar chemical

composition to the Japanese F-82H and JLF steels, to obtain an extensive understanding of,

and comprehensive database on, representative reduced activation steels. With this criterion,

the IEA modified F-82H and JLF-1 steels (large heats) were produced (27).

As a result of the investigations performed on the reduced activation ferritic/martensitic

steels, a primary candidate 9CrWVTa alloy, denominated EUROFER'97, was specified and

produced for the European DEMO breeding blanket concepts. Its qualification as a structural

material is being investigated nowadays and includes the characterization of conventional

properties and the determination of the irradiation behaviour up to damage levels of 70 dpa.

Low activation steels have a fully austenite structure when are austenizated in the temperature

range from 850°C to 1200°C. Austenite phase transforms to martensite phase during air

cooling or rapid cooling (quenching) to ambient temperature, and then steels are tempered to

obtain a good combination of strength, ductility, and toughness. However, the use of these

materials during long-time at high temperatures (thermal ageing) can produce microstructural

changes (new precipitates, grain growth, segregation, etc.) which can significantly affect their

mechanical properties (tensile, charpy, fracture toughness, low cycle fatigue, etc.). For these

reasons, an exhaustive knowledge of the metallurgical characteristics of these steels before

and after thermal ageing is considered essential.

This report describes the metallurgical properties (microstructural and mechanical) of the

Eurofer'97 steel, on as-received condition (normalized at 980°C/27' + tempered at

760°C/90'/air-cooled) and after thermal ageing treatments to simulate the service conditions,

in order to evaluate its feasibility as possible structural material for fusion applications.

3

2. EXPERIMENTAL PROCEDURE

2.1 Material

The Eurofer'97 is an European reference material within the framework of the European

Fusion Development Agreement (EFDA)-Structural Materials. This alloy has been fabricated

in Europe according to the chemical composition specifications required for the reduced

activation ferritic/martensitic steels type 9CrWTaV (Table 2).

The Eurofer'97 was fabricated in different products forms: plates, tubing, forging and wires.

Plates of 14 and 25 mm thickness were received at Ciemat for their study. These plates were

from the Heat E83698 (Plates3/8 and 3/9 of 14 mm thickness) and the Heat E83694 (Plate

1/13 of 25 mm thickness), as can be seen in Fig.1. All plates were supplied in the normalized

(980°C/27') plus tempered (760°C/90'/ air-cooled) condition.

In order to study the microstructural changes after long-term simulated service at high

temperatures, and their influence in the mechanical properties, the Eurofer'97 was thermally

aged at 400°C, 500°C and 600°C during 1000, 5000 and 10000 h.

2.2 Microstructural characterisation

The microstructural characterisation was performed along the three spatial orientations in

order to check the homogeneity of the plates. Microstructure was characterised by optical

microscopy (O.M) and scanning electron microscopy (SEM/EDS). In addition, TEM

investigations of thin foils (all material conditions) and carbon extraction replicas

(500°C/5000 h and 600°C/1000 h) were performed in a 200 kV JEOL transmission electron

microscopy equipped with and X-ray energy dispersive spectrometer (EDS). Thin foils were

prepared from 3 mm disks, grounded to a thickness of about 0.13 mm and finally

electropolished in an electrolyte containing 80% methanol and 20% sulphuric acid at -10°C.

Carbon extraction replicas for TEM studies were prepared from the metallographic samples

etched in a solution of Marble diluted. The specimens were then vacuum coated with a thin

carbon film. Re-etching in the same solution was then done to dissolve the matrix and extract

the precipitates from the surface. Carbon extraction replicas were washed in water and

mounted on copper grids for their examination.

On the other hand, phase extraction were performed by anodic dissolution of the matrix in

10% chloride acid-methanol containing 1% of tartaric acid under a current density of 0.1

4

A/cm2. The identification of second phase precipitates in the extracted residues was

performed by EDS (SEM) and by X-ray diffraction (XRD).

2.3 Mechanical properties

Vickers hardness measurements were carried out on as-received material and all aged

materials along the three spatial orientations. Tensile tests were also performed in all studied

material conditions using cylindrical specimens, machined from one of the 14 mm plates

parallely to the rolling direction, with 5 mm diameter and 25 mm of gauge length (Fig. 2).

The specimens were tested at room temperature and at the same temperature as that of the

ageing treatment.

Charpy impact tests were conducted on as-received and thermally aged materials at 500°C

and 600°C during 5000 and 10000 h, using V-notched specimens fabricated from one of the

14 mm plates with T-L orientation according to ASTM E-23. These tests were performed in

the temperature range from -120°C to 90°C in order to obtain the full impact curve. The

Ductile-Brittle-Transition-Temperature (DBTT) was estimated as 50% brittle fracture mode.

Low cycle fatigue tests were carried out on as-received and aged material at 500°C/5000 h

using round bar type specimens, with a gauge section of 8.8 mm in diameter and 25 mm

length (Fig. 3), manufactured from the 25 mm plate. The longitudinal direction of the

specimens was parallel to the rolling direction. The tests were carried out at 500°C in air.

Total strain ranges were 0.4, 0.7, 1 and 1.5 %, and the axial strain rate was of 5.10-3s-1.

Fracture toughness tests on as-received and aged material at 600°C/10000 h were performed

in the transition region in order to determine KJC values, with 1/2 TCT (1/2” thickness)

specimens (Fig. 4) machined from one of the 14 mm plates following the ASTM standard E

1921-02.

Creep tests are being carried out in the temperature range from 450°C to 650°C at different

loads, from 370 MPa to 50 MPa, using thread-head specimens of 5 mm-diameter x 25 mm of

gauge length (Fig. 2) fabricated from one of the plates of 14 mm.

5

3. RESULTS

3.1 Chemical composition

Chemical analyses to verify the composition of Eurofer'97 were performed at Ciemat. Table 3

summarises the results of these analyses together with the manufacturer analyses for

comparison. Both analyses are similar and agree quite well the chemical composition

specifications for this alloy (see table 2). Only a slight increase of the Ta concentration with

respect to the specifications has been observed, this difference being most significant in the

manufacturer analysis than in the Ciemat analysis.

3.2 Microstructural characterisation

Optical microscopy and scanning electron microscopy

The Eurofer'97 is a fully martensitic steel, free of δ-ferrite, with lath-shaped martensite

subgrains. Fig. 5 shows representative optical and SEM micrographs of this steel in the as-

received state. As it can be seen, the Eurofer'97 steel presents a fine structure with a prior

austenite grain size in the ASTM range 10-11.5 (6.7-11µm) for the 14 mm plate and ASTM

10.5-11 (8-9.4µm) for the 25 mm thickness plate. The tempering treatment produced large

amounts of carbide precipitation, distributed preferentially along grain and lath boundaries but

precipitates appear also in the bulk of the martensite laths.

Optical observations in the as-received state also revealed the presence of inclusions

randomly distributed and with spherical shape in the majority of the cases (Fig. 6), although

some long shaped inclusions were also detected. EDS analyses showed several kinds of

inclusion compositions, including MnS, Ta-rich inclusions and a few more complex

inclusions, possibly oxide or spinel-type, containing variable amounts of Al, Fe, Cr, Ti and V

(Fig. 7). The size of the inclusions was in the range 1.2-1.7 µm for MnS, 0.6-1.4 µm for Ta

rich inclusions, and oxide or spinel-type inclusions of approximately 4 µm were detected.

Starting from the as-received state, the Eurofer'97 steel was aged at 400°C, 500°C and 600°C

during 1000, 5000 and 10000 h. As it can be seen in table 4 the prior austenite grain size

remains stable at the heat treatments studied. Representative SEM micrographs of the aged

materials are shown in Figs. 8-10. No significant microstructural changes (prior austenite

grain size, morphology, size and distribution of the second phase precipitates) were observed

by optical and scanning electron microscopy in any of the aged materials.

6

Phase extraction and X-Ray Diffraction

Phase extraction was carried out on all material conditions studied. The results showed that

the amount of extracted residue does not increase significantly in aged materials compared to

the as-received state. In all conditions the amount of residue was in the range from 2.2% to

2.9%, Fig. 11. This suggests that no any significant new nucleation or growth of precipitates

have occurred during ageing.

The EDS (SEM) analyses performed in the extracted residues, Fig. 11, indicated the presence

of Cr, Fe, V, W, Ta and Ti. Except Ti, all these elements were detected in the as-received

condition. Figs. 12 and 13 show the alloying elements variation in the extracted residue in

function of the ageing time for each temperature and in function of the temperature for each

ageing time respectively. It is seen in these figures that the V, W, Ta and Ti contents remain

practically constant in all material conditions. The most pronounced variation is in the case of

Cr and Fe at 600°C/5000 h. The Cr content increases from 66.6 at % (as-received condition)

to 74.8 at % (600°C/5000 h) and the Fe concentration decreases from 24.6 at % (as-received

condition) to 16 at % (600°C/5000 h). Similar tendency of Cr and Fe level variation is also

detected in the materials aged at 500°C and 600°C for 10000 h, although less pronounced.

The X-ray diffraction patterns showed the same types of precipitates independently of the

material state, that is M23C6 as predominant carbide and precipitates with the same structure

than TaC. As an example, the X-ray diffraction patterns of as-received and aged materials at

400°C, 500°C and 600°C for 10000 h are illustrated in Figs. 14-17.

Transmission electron microscopy (TEM)

Fig.18a shows the general microstructure of Eurofer'97 at relatively low magnification. It is

characterised by prior austenite grain boundaries and martensite laths of 0.5±0.2 µm wide.

Regarding the secondary phases, two different types of precipitates (Fig. 18b) have been

identified by EDS analyses on carbon extraction replicas: Cr rich precipitates and Ta or V rich

precipitates. The main precipitation consists of Cr rich precipitates of variable size, from

approximately 25 nm to 200 nm (Fig. 19), located preferentially along grain and lath subgrain

boundaries but there were also some of them precipitated inside the subgrains. These particles

were identified as M23C6 type according to the results of XRD and electron diffraction. Their

morphology varied from globular to plates or to irregular geometrical shapes, and their typical

atomic concentration, obtained from extraction replicas, was 66±1 Cr / 31±1 Fe /1.9±0.2W,

7

though in some analysis V was found to replace W. Furthermore, in other analysis V was also

detected in addition to Cr, Fe and W.

In the as-received condition other type of precipitates, MX type, rich in Ta or V with smaller

size than M23C6, ranging from ~8 to 40 nm (Fig.19), were identified. They were mainly

located inside the subgrains. Three types of MX morphologies have been detected in the steel

after tempering (Fig. 20). Type I is a spherical Ta or V rich MX, this type being the most

numerous. Type II is a fine V-rich precipitate with a plate shape, and Type III represents a

specific morphology that is formed by secondary V precipitation at expense of the Ta rich

particles, so called V-wing.

EDS analyses performed on carbon extraction replicas in the as-received cond ition showed

that Ta rich precipitates contained about 60-80 at% Ta and 20 at% V. In the case of V rich

precipitates, V concentration was ∼70 at% and Ta ∼15 at%. Precipitates with equal amounts

of Ta and V (∼45 at%) were also observed. In addition, Fe and Cr were generally detected in

these analyses. These observations confirm the phase extraction results where a high V

content appeared, and also Ta, indicating the existence of Ta and V rich precipitates.

After the thermal ageing treatments investigated in this work, the Eurofer'97 steel showed

similar microstructural characteristics to those observed in the as-received condition. The

martensite laths wide values were practically the same for all the studied conditions, as it can

be seen in Fig. 21. Only a slight scattering in the data is observed for the aged material at

500°C/10000 h. The only different TEM feature was detected in materials aged at 500°C and

600°C for 10000 h. In aged material at 500°C/10000 h the occasional appearance of equiaxed

grains (Fig. 22) was found, characteristic that was also observed after ageing at

600°C/10000h, probably as a consequence of a recrystallization process in both cases. In the

majority of the cases, these grains were decorated with big M23C6 carbides (up to 350 nm) in

their boundaries and they were free or contained only a few dislocations and precipitates

inside them. In addition, TEM examination of a sample aged at 600°C/10000 h also revealed,

in some zones, the formation of the sub-grain structure replacing the martensite laths (Fig. 23)

and often with a polygonal shape. These recovery areas (sub-grain structure) were detected in

general close to the prior austenite grain boundaries.

The effects of ageing on the second phase precipitates were studied in carbon extraction

replicas of samples aged at 500°C/5000 h and at 600°C/1000 h. After these thermal ageing

8

treatments the morphology of M23C6 and MX precipitates was the same than in the as-

received condition. However, a slight particle growth was observed, more noticeable in the

case of M23C6, as it can be seen in Fig. 24. The size distribution of the second phases on each

material condition is shown in Fig. 25. The range of M23C6 particle diameter varied from ~ 40

to 260 nm for 500°C/5000 h, and from ~ 40 to 300 nm for 600°C/1000 h, while in the as-

received condition the range of particle diameter was ~ 25 to 210 nm. In the case of MX

precipitates the range size was ~ 10-60 nm and ~ 10-100 nm for 500°C/5000 h and

600°C/1000 h, respectively. The MX particles in the as-received condition exhibited lower

range size (~ 8-40 nm).

The EDS analyses performed on the precipitates also revealed small changes in their chemical

composition. The M23C6 presents at 500°C/5000h similar composition to that found on the as-

received condition (66±1 Cr / 31±1 Fe /1.9±0.2W at%). However, in the case of the material

aged at 600°C/1000 h, this type of carbide exhibited higher Cr and W concentrations (~ 71

at% and ~3.5 at% respectively) and lower Fe content (~ 26 at%) than in the other conditions

(as-received and aged material at 500°C/5000 h). Regarding the MX type precipitates, the

most important finding was that after the ageing treatments at 500°C and 600°C, Ta pure

particles were found that were not seen in the as-received condition. In addition, a few MX

particles with Ti were also found.

3.3 Mechanical properties

Hardness measurements

The HV30 hardness values of Eurofer'97 are summarised in table 5. Each value represents the

mean value of the three spatial orientations. In the as-received condition, the Eurofer'97 steel

exhibits hardness values of approximately 210±3, hardness level that remains quite stable

after the ageing treatments studied. These results agree quite well with the TEM

investigations and phase extraction results where no new precipitation or significant growth of

martensite laths were detected.

Tensile properties

Tensile strength, yield strength, total elongation and reduction of area values of Eurofer’97 in

the material conditions studied in this work are plotted in Fig. 26. The steel on as-received

condition exhibits ultimate tensile strength and yield strength values at room temperature of

662 and 530 MPa, respectively. After the ageing treatments investigated, the tensile

characteristics of the Eurofer'97 do not show any variation. There are no differences in the

9

ultimate tensile strength and yield strength values for each temperature in function of the

ageing time.

Charpy properties

The impact curves of the Eurofer'97 steel on as-received state and after the thermal ageing

treatments at 500°C and 600°C for 5000 and 10000 h are presented in Fig. 27. The Eurofer'97

steel exhibits in the as-received condition an upper shelf energy value of 266 J and a DBTT of

-51°C. The USE remains practically constant after ageing at 500°C and 600°C up to 10000 h.

The most significant variation of the DBTT takes place in the material aged at 600°C during

10000 h, with a DBTT increase of approximately 23°C with respect to the as-received state.

Fracture toughness

Fracture toughness tests were performed in the Eurofer'97 steel in the as-received condition

and in the material aged at 600°C/10000 h, in order to determine To reference temperature

following the Master Curve approach that was developed by VTT (28) and that is based on

the analysis of cleavage fracture data. Following this approach, dependence of fracture

toughness with temperature can be described by:

KJC (med) = 30 +70 exp (0.019(T-To))

Where:

KJC (med) is the median value of the Weibull distribution that describes the scatter of fracture

toughness results. To is the reference temperature at which KJC(med) has a value of 100

MPa√m for a 1" specimen thickness.

To reference temperature for Eurofer'97 was determined following the multi-temperature

technique inc luded in ASTM E1921-02. Experimental fracture toughness results and Master

Curve for the as-received material can be seen in Fig. 28. The scatter of the data obtained in

the fracture toughness tests is the normal found in this type of tests. The T0 value obtained for

the Eurofer'97 steel in the as-received condition was of -129°C (multi- temperature).

Previously (29), T0 was calculated in the Eurofer'97 steel on as-received condition according

to ASTM 1921 May 2000 Draft using single temperature (T0 = -126°C) and multi-

temperature (T0 = -132°C) technique.

10

Results of fracture toughness tests on aged material at 600°C/10000 h can be seen in Fig. 29

In this case T0 value is –122 ºC (multi- temperature). The difference of 7 ºC between the two

tested conditions can not assure that the material has lower fracture toughness in the aged

condition than in the as-received condition because it could be attributed to the scatter of the

results. Anyway, this small difference should not be discarded taking into account that, in

Charpy tests, DBTT showed an increase of 23 ºC in this aged material condition compared to

the as-received condition.

Low Cycle Fatigue properties

The strain- life curves of the Eurofer'97 on as-received state and after ageing at 500°C for

5000 h, both tested in air at 500°C are plotted in Fig. 30. The values corresponding to 0.4%

total strain range for the material aged at 500°C/5000 h have not been included because the

failure took place in the threads. The results indicate that the fatigue behaviour of Eurofer'97

steel after thermal ageing at 500°C/5000 h is quite similar to the as-received steel for total

strain ranges of 1.5% and 1%. Nevertheless, the number of cycles to failure at 0.7% total

strain range is slightly lower than in the as-received state. To be sure of the influence of the

studied ageing treatment in the fatigue life time of the Eurofer'97, more tests at 0.4% total

strain ranges should be necessary.

The strain- life curve corresponding to the as-received condition was also represented

following the Manson-Coffin equation:

∆εt = ∆εp + ∆εe = CpNf(-kp) + CeNf(-ke)

where ∆εp and ∆εe are the plastic strain and the elastic strain ranges, respectively. Cp, Ce, kp and ke are material constants. According to this equation, the relations between plastic and

elastic strain ranges and cycles to failure are shown in Fig. 31.

The change in the tensile peak stress during fatigue- life is shown in Figs. 32 and 33. In both

material conditions, the stresses decreased continuously up to the final failure for 0.7%, 1%

and 1.5% total strain ranges. However, apparently in the tests performed in the as-received

condition at 0.4% total strain range, the peak stress increased up to cycle 10 and then

decreased gradually up to the final failure. These results could suggest that the Eurofer'97

experiments hardening at this strain. Nevertheless, in order to clarify this observation a

detailed evaluation of the tests at 0.4% was performed. The raw data showed that the strain

11

range was lower than the one required (0.4%) up to cycle 10, then the total strain range started

to be stable at 0.4%. Consequently the peak stress increased from the beginning up to cycle 10

because it was necessary higher stress to reach the 0.4% total strain range. Therefore, the

Eurofer'97 steel experiments continuous cyclic softening for the higher strain ranges (0.7%,

1% and 1.5%) and further tests at 0.4% should be necessary to verify this trend at this strain

range.

Creep properties

Fig. 34 shows some representative creep and creep rate curves of Eurofer'97 for each test

temperature. The creep rate curves consist of a primary stage where the creep rate decreases

with time, a secondary region also denominated a minimum creep rate, and a tertiary or

accelerated creep region, where the creep rate increases with time.

The stress rupture data for the testing temperatures are plotted in Fig. 35. The Eurofer'97

exhibits adequate creep rupture strength levels in the range of temperatures and loads tested,

except at 500°C and 200MPa (15470 h) whose results seem to suggest some deterioration of

the creep rupture properties.

There are different methods to evaluate the creep rupture strength of one steel for long times.

One of them widely used is the Larson-Miller parameter (30), Fig. 36. This parameter

assumes that temperature and time can be interchanged, that is long times at low temperatures

are equivalent to shorter times at higher temperatures, provided no important microstructural

changes occur during the creep. However, in the case of Eurofe r'97 not all the data (specimen

tested at 500°C, 200MPa during 15470 h) seem to fit the creep master curve with accuracy.

Another assessment of the data is using the Norton equation (εmin = kσn), which relates the

variation of the minimum creep rate with the applied stress. This equation describes the creep

deformation characteristics because a change of the Norton stress exponent (n) is indicative of

a change in deformation mechanism. The n exponent of the Eurofer'97 steel was determined

for each test temperature (Table 6). The examination of the results, Fig. 37, seems to confirm

that at 500°C and the lowest stress a creep behaviour change has occurred. This could be

indicating that a microstructural change is taking place during the creep test, but additional

results would be necessary to determine the n exponent at lower stresses and to corroborate

that the Eurofer'97 steel at 500°C exhibits two regions with different deformation mechanism.

In addition, it is possible that at 550°C and the lowest stresses (160MPa) a change in the creep

behaviour might also been occurring.

12

In order to know the microstructural development during the creep tests, some representative

samples were selected for microstructural investigations (phase extraction, X-ray diffraction

and TEM/EDS). At short rupture times (9000h), in addition to the M23C6 and MX

precipitates, the microstructural investigations performed by TEM in thin foils revealed the

presence of new particles. This new phase (Fig. 39a) tentatively identified as M6C type

(Fe3W3C) according to the X-ray diffraction results (Fig. 39b) is mainly precipitated at the

sub-grain boundaries and in the majority of the cases is associated to the M23C6 carbides.

Typical EDS spectra of these M6C and M23C6 carbides are shown in Fig. 40 a and b

respectively. As it can be seen, the analyses show high Fe and W concentrations in the M6C

type particles.

13

4. DISCUSSION

The Eurofer'97 steel is essentially a low carbon steel with similar composition to other

reduced activation martensitic steels, such as Optifer, JLF-1 and LA series (31). From the

point of view of reduced activation criteria and chemical composition, the Eurofer'97 agrees

quite well the specifications proposed for this alloy (32). It is well stablished that small

differences in the concentration of some alloying elements, such as Ta and V, have strong

influence on the microstructure of the alloy. The fine structure observed in the Eurofer'97 is

basically attributable to the Ta concentration in the steel (0.1 wt%), Ta content that has also

influence on the type of existing precipitates and on the mechanical properties as it will be

discussed later.

If the as-received Eurofer'97 steel (9Cr-1WVTa) is compared with the as-received reduced

activation ferritic/martensitic steel F-82H modified (8Cr-2WVTa) previously studied (33),

some microstructural differences can be observed related with the amount and type of

precipitates that are formed depending on the chemical composition of the alloy. Phase

extraction results, Fig. 41, show that the percentage of extracted residue was higher in the

Eurofer'97 than in the F-82H mod. steel besides of the differences found related to the

elements present in the extracted residues. The EDS (SEM) analyses performed on extracted

residues showed that, apart from Fe, Cr, W and V present in both steels, also Ta was detected

in the Eurofer'97. In addition, V was observed in higher concentrations in the Eurofer'97 than

in the F-82H mod. steel. In accordance with these results, TEM investigations confirmed the

existence of different type of precipitates in the Eurofer'97 steel: M23C6, MX rich in Ta and

MX rich in V. In contrast, in the F-82H mod. steel only M23C6 carbides were detected. This

could be due to the fact that in this steel the concentration of Ta (0.005wt%) and V (0.14wt%)

are low to form MX precipitates. In addition to the M23C6 and/or MX precipitates, other type

of precipitates have been observed in other reduced activation ferritic/martensitic steels

depending on the alloy composition, that is the case of the LA series (34) and the JLF-1 steels

(35). M2X particles rich in V were identified in Cr-rich LA4Ta steel

(11Cr/0.8W/0.1Ta/0.23V), and M6X Ta-rich particles and M4X3 V-rich precipitates were

detected in LA4Ta, LA12Ta (10Cr/0.8W/0.1Ta/0.3V), LA13Ta (9Cr/3W/0.1Ta/0.24V) and

JLF-1 (9Cr/2W/0.1Ta/0.2V) steels.

It is well known (35,36) that MX type particles are very useful for long term creep resistance

at elevated temperatures. The three types of MX morphologies identified in the Eurofer'97 in

this work are similar to those found by Yamada and co-workers (37) in experimental

14

9CrWNbV steels. These authors also reported that the formation of Type III-MX precipitates,

with its specific feature in form of V-wing, depends on the N concentration in the steel. The

investigations performed by Tamura and Yamada (36,37) seem to indicate that the minimum

N concentration necessary to form this MX type is around 0.02%, that is the same N content

of the Eurofer'97. Nevertheless, in the F-82H mod. this type of precipitate was not detected,

also in accordance with its N content (0.007%) that is lower than in the Eurofer'97 and in the

other mentioned 9Cr steels. Yamada also reports that Type III-MX is formed from Nb(C,N).

These particles, formed during the normalization treatment, have a region rich in V on the

surface and consequently they can act as nucleation sites during tempering, growing the VX

like a wing.

On the other hand, it is well recognized that the microstructural stability of one steel at high

temperature during long times (simulated service conditions) is the vital importance for its

mechanical behaviour. All microstructural transformations that could take place in the steel

during high temperature service will have a critical influence on its mechanical properties

and, consequently, even they would make questionable the use of this steel.

The high microstructural stability observed in this work by optical and scanning electron

microscopy in the Eurofer'97 steel after the investigated thermal treatments was confirmed by

TEM. The only difference detected by TEM was the presence of some equiaxed grains in the

material aged at 500°C and 600°C for 10000 h. These grains probably recrystallized during

the ageing treatments, because neither in the as-received state nor in the other studied material

conditions were found. In addition, the presence of some big M23C6 carbides at the equiaxed

grain boundaries seems to indicate a recrystallization process because it is well known (37)

that the mobility of lath interfaces increases as the growth and coarsening of the M23C6

carbides proceed, specially if these particles are principal responsible for the pinning of the

lath boundaries. Consequently, such growth and/or coarsening would be expected to promote

the tendency for the progressive breakdown of the martensite lath morphology during

prolonged exposure at high temperature, as it seems to have been the case of the Eurofer'97

aged at 500°C and 600°C for 10000 h. Partially recrystallized regions have also been

observed in the F-82H mod. steel in the as-received condition as well as after thermal ageing

at 600°C for 5000 h and at 550°C for 13500 h (33,34).

Regarding to the M23C6 and MX precipitates, it is clear from Fig. 24 that the ageing

treatments at 500°C/5000 h and at 600°C/1000 h could be considered equivalent thermal

15

ageing treatments from the point of view of particle growth. The M23C6 and MX presented

similar mean particle diameter values for both ageing conditions. However, the M23C6 carbides showed higher tendency to growth than the MX particles during ageing. J. Hald and

M. Hättestrand (39-41) have extensively reported the major stability of MX precipitates

during ageing or creep. On the other hand, each ageing treatment studied in this work had

different influence on the chemical composition of the precipitates, particularly in the case of

the M23C6 carbides. The Cr content increases and the Fe concentration decreases in aged

material at 600°C/1000 h while at 500°C/5000 h the composition remains constant, both

conditions compared with the as-received state. These observations have also been confirmed

by the phase extraction results where a higher Cr content increase was detected after the

ageing treatments at 600°C. Similar trend in the Cr and Fe carbide concentrations has been

observed by other authors (33, 42) as a result of ageing treatments at 550°C and 600°C for

long times.

Relating to the tensile properties, the Eurofer’97 steel presents similar levels of ultimate

tensile strength and yield stress than other reduced activation ferritic/martensitic steels (F-

82H, Optifer II and Batman), as can be appreciated in Fig. 42. The tensile characteristics of

the reference MANET II (11Cr/Mo/V/Nb/Ni) structural material for fusion applications are

also included for comparison. The tensile properties of this steel are considered as the

reference properties that must have reduced activation ferritic/martensitic steels. As can be

seen in the graph, at room temperature the higher values correspond to the Optifer Ia

(9Cr/1W/0.06Ta/0.26V) and the steel 91 (9Cr/1Mo/0.06Nb/0,25V). This behaviour is related

to the higher hardness level of these materials (Optifer Ia and steel 91) than the other alloys

after the tempering treatment (750°C/1-2h). However, above 550°C the differences between

the alloys are not significant (31).

The continuo composition optimisation of reduced activation ferritic/martensitic steels, in

order to improve their mechanical properties and, specially, the impact properties, has given

rise to obtain DBTT values of approximately –51°C for the Eurofer’97 steel, and even smaller

for the Optifer IV (around –60°C), against to DBTT values of –25°C in the case of the F-82H

steel and near to the room temperature for the steel 91 (Fig. 43). These DBTT differences are

related with the alloying elements concentration and the heat treatments applied. For example,

the better toughness of the Eurofer’97 steel than the one of F-82H steel can be attributable to

the W, Ta and V concentrations present in the alloys and the different normalization

treatment (980°C/27’ for the Eurofer’97 and 1040°C/37’ for the F-82H mod.). In

16

investigations performed by Abe and co-workers (43,44) the effect of these alloying elements

on the impact properties of simple Cr-W and of reduced activation CrWTaV steels were

systematically studied. They concluded that DBTT optimum values are obtained with W

concentration around 1 wt.%, as it is the case of the Eurofer'97, and when the steel presents a

fine microstructure due to small addition of alloying structure refinement elements, such as Ta

and V, also in similar concentrations to those present in the Eurofer’97 steel. As a

consequence of both factors (elements concentration and heat treatment), the grain size is bout

6.7-11µm in the Eurofer’97, while is approximately 65-75 µm in the F-82H mod. steel. On

the other hand, the higher DBTT of the steel 91 seems to be related with the austenization

treatment that in this case was 1040°C (31). This temperature seems to be very high to obtain

a small grain size and hence optimum fracture toughness, due to for these alloys it is

recommendable to lower end of the austenization range (950°C-1050°C) in order to avoid a

grain growth. The austenization treatment of the Optifer IV steel (900°C/30’) led to a smaller

grain size (∼9.5 µm) than the other alloys (31), fact that could be one of the principal reasons

for its low DBTT.

After the investigated ageing treatments, the mechanical properties of the Eurofer'97 steel do

not show any significant degradation in accordance with the high microstructural stability

observed. Only a slight increase of the DBTT, of about 23°C, was detected after ageing at

600°C/10000 h. These results agree quite well with the observations described by several

authors (45-47) indicating that no significant hardening-embrittlement is detected in 9Cr

martensitic steels thermally aged from 500°C to 650°C up to 10000 h. These studies show a

DBTT increase, from the as-received condition, of approximately 20°C at 500°C and 90°C at

650°C, but not any explanation related with microstructural changes is given. In this work,

taking into account the microstructural investigations performed in the aged Eurofer'97, the

slight DBTT increase might be related with the martensite transformation to subgrains after

long-term ageing. However, no references have been found in the literature about a possible

relation of these microstructural features with the impact properties.

Turning next to the fracture toughness there is some controversy with respect to the Master

Curve application to martensitic steels. Anyway, Master Curve is being applied to martensitic

steels with good results (48-55). G.R. Odette (52) assumes that Master Curve could be applied

to static loads with fracture in or near the Small Scale Yielding (SSY) regime for the

martensitic steel F82H. This assumption could be applied to Eurofer'97 fracture toughness

17

tests as they have been performed with static load and with a geometry (1/2TCT, a/W=0.5)

that would assure near SSY regime.

The Master Curve of the Eurofer'97 steel in the tested material conditions seems to show a

relatively sharp transition. Similar tendency of Master Curve was obtained by J.W. Rensman

in the Eurofer'97 steel (54), but in this case the T0 in the as-received condition was 29°C

higher than the T0 ( -129°C) calculated in Ciemat. It is worth pointing out that the specimens

used by J.W. Rensman were mechanized from a 25mm thickness plate while that Ciemat

specimens were fabricated from a 14mm thickness plate. The author in other work (55) also

obtains values of DBTT, measured by charpy tests in specimens fabricated from a 25 mm

thickness plate, about 25°C higher than the rest of the product forms of Eurofer'97 steel (8

mm and 14mm thickness plate and bar). This author indicates that the 25 mm plate has

impaired deformation capacity and toughness by microstructural origin on TEM-scale. Then,

the difference in T0 values between the two plates is not unexpected.

G. R. Odette et al. (53) plot the variation of the Eurofer’97 fracture toughness in function of

the specimen thickness at –142°C, as can be seen in Fig. 44. From this figure, the Eurofer’97

steel at –142 ºC give a fracture toughness of about 78 MPa√m with 1” thickness specimens.

Extrapolating the test temperature of –142°C in the Ciemat Master Curve (T0 = -129°C), Fig.

28, the Eurofer’97 presents fracture toughness values of approximately 84.7 MPa√m. The

results are very similar with only a little difference of 6.7 MPa√m. A comparison can also be

established with the F82H martensitic steel (56). This work is a compilation of fracture

toughness of the F-82H using different type and size of specimens (Fig. 45). The authors give

a value of T0 = -119.3 ºC for this material, which is slightly lower than the Eurofer’97 T0 value obtained in this work.

It is known that fatigue softening is a typical phenomenon for materials hardened by

precipitation, solid solution and/or martensitic transformation (57). Fatigue softening is

caused by the gradual elimination of obstacles, such as precipitates, foreign particles and

grain or lath boundaries, to the motion of dislocations. Although fatigue softening curves

usually have a saturated region, the curves obtained in this work for the as-received

Eurofer'97 steel tested at 500°C (Fig. 32 and 33) did not show any saturation. This behaviour

has also been observed in the F-82H steel (57). The authors attribute this behaviour to the

continuos change of the distribution and morphology of obstacles. On the other hand, it is

worth pointing out that the peak stress curve obtained at 0.4 % total strain range does not

18

mean that Eurofer'97 experiments cyclic strain hardening by the reasons previously

mentioned in the section of low cycle properties. The low cycle fatigue results are discussed

in a work of compilation and evaluation of Eurofer'97 properties performed by Tavassoli (58).

The author, considering only the stress curves, reports that the alloy exhibits softening at high

strain ranges and hardening at low strain ranges. However, as it was discussed in the fatigue

section of this work, it is necessary to take also into account the strain range values obtained

during the tests to be sure that the Eurofer'97 experiments hardening at 0.4% total strain

ranges. On the other hand, it is difficult to evaluate the fatigue strength of the Eurofer'97 aged

at 500°C/5000 h basically for two reasons. One of them is the fact that more tests at 0.4%

total strain range are necessary in order to obtain a good curve fit and, consequently, be able

to compare the fatigue behaviour of the aged Eurofer'97 with the as-received state. The other

reason is related with the lack of low cycle fatigue data published for reduced activation steels

after ageing. Unfortunately no references have been found in the literature about low cycle

fatigue tests of aged materials.

Concerning the creep properties, the Eurofer'97 exhibits adequate creep rupture strength level

at short creep rupture tests similar to other reduced activation ferritic/martensitic steels, such

as the F-82H mod. steel (38). However, at long testing times (> 9000 h), the results available

up to now at 500°C and 550°C suggest that a change in the deformation mechanism could be

taking place. The precipitation of M6C type precipitate, with high W content, could be the

reason of the creep strength degradation of Eurofer'97 steel for long testing times. It is well

known that W is a solid solution strengthening element. The formation during the creep tests

of new precipitates rich in W, as it seems to be the case of Eurofer'97, produces a depletion of

W in the matrix. The result is a loss of solid solution hardening that contributes to the

deterioration of the creep properties. Similar precipitation of M6C carbides during creep at

550°C and long test times has been observed in several casts of 12CrMoVNb steels (59,60),

in the modified 9Cr-1Mo steel (61) and in a 1Cr-0.5Mo steel (62). The authors indicate that

the precipitation of M6C phase plays an important role in the abrupt lowering of creep rupture

strength of these steels, these observations being in good agreement with the results obtained

in this work for the Eurofer'97 steel.

19

5. SUMMARY AND CONCLUSIONS

- The Eurofer'97 on as-received condition is a fully martensitic steel free of δ-ferrite,

with several type of inclusions heterogeneously distributed, and with a fine grain size.

In addition, the material presents a high degree of homogeneity along the three spatial

directions with respect to the grain size and hardness values. Two kinds of precipitates

with different morphologies, namely M23C6 and MX (Ta or V rich), have been

detected in the material. After the ageing treatments investigated in this work, from

400°C to 600°C up to 10000 h, the steel did not exhibit significant microstructural

changes.

- The Eurofer'97 on as-received condition exhibited adequate strength and ductility

levels, comparable to other reduced activation ferritic/martensitic steels. The

Eurofer'97 steel showed, for each test temperature, similar values of ultimate tensile

strength and 0.2% proof stress regardless of the ageing condition.

- No significant change in the USE was observed after the studied ageing treatments.

Only a slight increase of the DBTT, of approximately 23°C, was observed after ageing

at 600°C/10000 h.

- Fracture toughness of Eurofer'97 steel follows the Master Curve with T0 values of -

129°C for the as-received condition and -122°C for the material aged at

600°C/10000h.

- The low cycle properties of Eurofer'97 tested at 0.7%, 1 % and 1.5% total strain range

do not seem to degrade significantly after ageing at 500°C/5000 h. Softening of

Eurofer'97, without showing a saturated region was observed in the fatigue curves at

500°C, independently of the material condition and the total strain ranges. Further

tests at 0.4% total strain range should be performed to evaluate entirely the fatigue

behaviour of the Eurofer'97 after ageing at 500°C/5000 h.

- The Eurofer'97 steel presents adequate creep rupture level for short-term tests. Taking

into account the microstructural investigations, the precipitation of M6C type particles

seems to be the principal reason for the creep properties degradation at 500ºC of the

Eurofer'97 steel for long testing times.

20

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60. V. Foldyna, Z. Kubon, V. Vodarek and J. Purmensky. Proceedings. of the 3rd Conference

on Advanced in Material Technology for Fossil Power Plants. Edited by R. Viswanathan,

W.T. Baker, and J.D. Parker. IOM Book 0770, (2001), 89.

61. P. Anderson, T. Bellgardt and F.L. Jones. Mat. Sci. Tech., 19, (2003), 207.

62. J. Dobrzanski and A. Hernas. J. Mat. Proces. Tech., 53, (1995), 101.

23

USA CEC JAPAN

2-9Cr-V 9Cr-W-V-Ta-N 2-15Cr-W

2-9Cr-W 12Cr-W-V-Ta-N 2-3Cr-W-V-Ta

2-12Cr-W-V 9-10Cr-W-V-Ta-Ti-Ce 7-9Cr-W-V-ta

9Cr-W-Mn 9Cr-W-V-Mn-Ti 11Cr-W-V-Ta

9Cr-V-Mn

12Cr-W-Mn

12Cr-V-Mn

Table 1: Basic composition of reduced-activation ferritic/martensitic steels.

24

ELEMENT MIN. Value (Wt%) MAX. Value (Wt

%) Remarks Target

C 0.090 0.120 0.11

Mn 0.20 0.60 0.4

P 0.005

S 0.005

Si 0.050

Ni 0.005 ALAP

Cr 8.50 9.50 9

Mo 0.005 ALAP

V 0.15 0.25

Ta 0.05 0.09

W 1.0 2.0

Ti 0.01

Cu 0.005 ALAP

Nb 0.001 ALAP

Al 0.01 ALAP

N 0.015 0.045

B 0.001 ALAP

Co 0.005 ALAP

As+Sn+Sb+Zr 0.05 Target

O 0.01

ALAP: As Low As Possible.

Table 2: Chemical composition specifications of Eurofer'97.

25

ELEMENTS HEAT E83698

PLATES 3/8 AND 3/9 (14 mm)

HEAT E83694

PLATE 1/13 (25 mm)

Manufacturer

(wt%)

Ciemat (wt%)

Manufacturer (wt%)

Ciemat (wt%)

C 0.11 0.11 0.10 0.10

Cr 8.82 8.7 8.87 8.7

W 1.09 1 1.15 1.1

Ta 0.13 0.10 0.14 0.11

V 0,20 0.19 0.20 0.19

Mn 0.47 0.44 0.45 0.42

Si 0.040 0.041 0.050 0.05

P 0.005 - 0.005 -

S 0.004 0.004 0.004 0.004

Mo

26

PAG As-

received 400°C/1000 h 400°C/5000 h 400°C/10000 h 500°C/1000 h 500°C/5000 h 500°C/10000 h 600°C/1000 h

600°C/5000 h 600°C/10000h

(ASTM) 10-11.5 11-11.5 10.5-11 10.5-11 11-11.5 10.5-11 10-11 10-11 10-11 10-11

µm 6.7-11 6.7-8 8-9.4 8-9.4 6.7-8 8-9.4 8-11 8-11 8-11 8-11

PAG: Prior Austenite Grain Size.

Table 4: Prior austenite grain size of Eurofer'97 in the as-received condition and after ageing.

As-

received 400°C/1000 h 400°C/5000 h 400°C/10000 h 500°C/1000 h 500°C/5000 h 500°C/10000 h 600°C/1000 h 600°C/5000 h 600°C/10000 h

HV30 210±3 212±4 213±1 206±3 214±3 211±3 209±2 207±1 201±3 208±2

Table 5: Vickers hardness values of Eurofer'97 in the as-received condition and after ageing.

27

Temperature n Loads 450°C 21 370-300 500°C 22 290-250 550°C 16 230-170 600°C 10 150-100 650°C 7 110-60

Table 6: Calculated "n" exponent for each test temperature.

28

Figure 1: Eurofer'97 distribution.

3260

1000 1000

500 500 500 500 500 500 260

400

800

CEA-S CEA-S CEA-SCEA-S ENEA-Br ENEA-Br

CIE

MA

T

CEA-G CEA-G

ENEA-Ca ENEA-Ca

ENEA-Ca ENEA-Ca

Res.

Res.

400

400

400

1200

1/3 1/4 1/5 1/6 1/7 1/8 1/13

1/9 1/10

1/11 1/12

1/14

1/15

Heat E83694Plate Nº 1t = 25 mm

1/1 1/2

2570

1000

500 500 600 600 37020

080

0

ENEA-Br NRG CIEMAT

CEA-G

FZK-IRS

Res.

Res.

200

400

400

1000

3/6 3/7 3/8 3/9 3/10

3/5

3/11

3/12

Heat E83698Plate Nº 3t = 14 mm

3/1 3/3

FZK-IRS

CIEMAT Res.

FZK-IRS

3/2 3/4

FZK-IRS

29

Figure 2: Tensile and creep specimen geometry.

30

Figure 3: Low Cycle Fatigue specimen.

31

Figure 4: Fracture toughness specimen (1/2TCT).

60º r=0.1

25.4 0.1±

31.75± 0.2

12.24± 0.1

30.4

8± 0

.2

φ= 6.35

1.5

12.7± 0.1

2.54± 0.1

2.545.349.53± 0.1

1.6

1.6

+0.1-0

32

(a)

(b)

Figure 5: a) Optical, b) SEM micrographs of Eurofer'97 in the as-received condition.

33

Figure 6: Optical micrograph of the inclusions.

Figure 7: EDS of the different inclusions type present in the Eurofer'97 in the as-received condition.

Fe

S CrMn

Fe MnS

Ti

Ta

VCr

Fe

Ta

Ta-rich Spinel-TypeAl

Fe STi

VCr

MnFe

Spinel-Type

34

Figure 8: SEM micrographs of Eurofer'97 aged at 400°C: a) 1000h, b) 5000h and c) 10000h.

(a) (b)

(c)

35

Figure 9: SEM micrographs of Eurofer'97 aged at 500°C: a) 1000h, b) 5000h and c) 10000h.

(b) (a)

(c)

36

(c)

Figure 10: SEM micrographs of Eurofer'97 aged at 600°C: a) 1000h, b) 5000h and c)

10000h.

(b) (a)

37

Figure 11: Phase extraction results of the Eurofer'97.

0

10

20

30

40

50

60

70

80

90

100

Cr Fe V W Ta Ti

2,5* 2,6* 2,9* 2,9* 2,4* 2,5* 2,2* 2,4* 2,5* 2,4*

* Percentage of extracted residue (wt%)

AR 400°C1000h400°C5000h

400°C10000h

500°C1000h

500°C5000h

500°C10000h

600°C1000h

600°C5000h

600°C10000h

Material Condition

Con

cent

ratio

n (a

t%)

38

Figure 12: Variation of the alloying elements in the extracted residue of aged Eurofer'97.

Influence of the ageing time.

As-received 600°C/1000h 600°C/5000h 600°C/10000h0

20

40

60

80

Con

cent

ratio

n (a

t %)

Material Condition

As-received 500°C/1000h 500°C/5000h 500°C/10000h0

20

40

60

80As-received 400°C/1000h 400°C/5000h 400°C/10000h

0

20

40

60

80 Cr Fe V W Ta Ti

Con

cent

ratio

n (a

t %)

Con

cent

ratio

n (a

t %)

Cr

Fe

Cr

Fe

Fe

Cr

39

Figure 13: Variation of the alloying elements in the extracted residue of aged Eurofer'97.

Influence of the ageing temperature.

As-received 400°C/10000h500°C/10000h600°C/10000h0

20

40

60

80

Con

cent

ratio

n (a

t %)

Con

cent

ratio

n (a

t %)

Con

cent

ratio

n (a

t %)

Material condition

Cr Fe V W Ta Ti

As-received 400°C/5000h 500°C/5000h 600°C/5000h0

20

40

60

80As-received 400°C/1000h 500°C/1000h 600°C/1000h

0

20

40

60

80

Cr

Fe

Fe

Cr

Cr

Fe

40

Figure 14: X-Ray Diffraction pattern of Eurofer'97 Figure 16: X-ray diffraction pattern of Eurofer'97 in the as-received condition. aged at 500°C/10000h

Figure 15: X-Ray Diffraction pattern of Eurofer'97 Figure 17: X-ray diffraction pattern of Eurofer'97 aged at 400°C/10000h aged at 600°C/1000h.

0 2 0 4 0 6 0 8 0 1 0 00

1 0 0 0

2 0 0 0

3 0 0 0

4 0 0 0

5 0 0 0

6 0 0 0

E u r o f e r ' 9 7 A s - r e c e i v e d

M23

C6

M23

C6

M 2 3 C 6

M23

C6

M2 3

C6

TaC

T a CTaC

TaC

Cou

nts

2 T h e t a

0 2 0 4 0 6 0 8 0 1 0 00

1 0 0 0

2 0 0 0

3 0 0 0

4 0 0 0

5 0 0 0

M23

C6

M23

C6

M2 3

C6

M23

C6

M 2 3 C 6

E u r o f e r ' 9 7 A g e d a t 4 0 0 ° C / 1 0 0 0 0 h

T a C

TaCT

aC

TaC

Cou

nts

2 T h e t a

0 2 0 4 0 6 0 8 0 1 0 00

1 0 0 0

2 0 0 0

3 0 0 0

4 0 0 0

5 0 0 0

6 0 0 0

7 0 0 0

8 0 0 0

M23

C6

M23

C6

M23

C6

M 23 C 6

M 2 3 C 6

Ta

CTaC

TaC

T a C

Cou

nts

2 T h e t a

E u r o f e r ' 9 7 A g e d a t 5 0 0 ° C / 1 0 0 0 0 h

0 2 0 4 0 6 0 8 0 1 0 00

5 0 0

1 0 0 0

1 5 0 0

2 0 0 0

2 5 0 0

3 0 0 0

3 5 0 0

TaC

T a CTa

C

TaC

M23

C6

M2

3C6

M 2 3 C 6

M23

C6

M 2 3 C 6Cou

nts

2 T h e t a

E u r o f e r ' 9 7 A g e d a t 6 0 0 ° C / 1 0 0 0 0 h

41

Figure 18: a)General microstructure by TEM of the Eurofer'97 in the as-received condition

and b) M23C6 and MX precipitates of the Eurofer'97 steel in the as-received condition.

M23C6

MX

(a)

(a)

42

Figure 19: Size distribution of Eurofer'97 precipitates in the as-received condition.

Figure 20: MX precipitate morphologies of the Eurofer'97 in the as-received condition.

0 5 10 15 20 25 30 35 40 450

5

10

15

20

25

MX

Particle diameter (nm)

0 20 40 60 80 100 120 140 160 180 200 2200

5

10

15

20

25

30

35

Fre

cuen

cy(%

)M

23C

6

Fre

cuen

cy(%

)

VTa

V-WingPlate

Spherical

43

Figute 21: Mean laths wide of Eurofer'97 as a function of the material condition.

As-received 500°C/5000h 500°C/10000h 600°C/1000h 600°C/5000h 600°C/10000h0,0

0,1

0,2

0,3

0,4

0,5

0,6

0,7

0,8

0,9

1,0

Mean laths wide ± sd

Material Condition

Lath

s Wid

e (µ

m)

44

(a)

(b)

Figure 22: TEM micrographs of the Eurofer'97 steel aged at 500°C for 10000 h: a) equiaxed

grains and b) incipient formation of the equiaxed grains.

45

Figure 23: Sub-grain structure in Eurofer'97 steel aged at 600°C for 10000 h.

46

Figure 24: Particle mean diameter of M23C6 and MX precipitates in the as-received condition

and after ageing at 500°C/5000 h and at 600°C/1000 h.

Figure 25: Particle size distribution of Eurofer'97 in the as-received state and after thermal

ageing at 500°C/5000 h and at 600°C/1000 h.

As-received 500°C/5000h 600°C/1000h0

25

50

75

100

125

150

175

200

Par

ticle

dia

met

er (n

m)

Material condition

Mean particle diameter of M23C6 ± sd Mean particle diameter of MX ± sd

0 20 40 60 80 1000

5

10

15

20

25

30

MX

As-received 500°C/5000 h 600°C/1000 h

Fre

cuen

cy (

%)

Particle diameter (nm)

0 50 100 150 200 250 3000

5

10

15

20

25

30

(b)

(a)M

23C

6

Fre

cuen

cy (%

)

As-received 500°C/5000 h 600°C/1000 h

47

Figure 26: Tensile Tests of Eurofer'97 (RT = Room Temperature).

Ageing Tª = Test Tª

0 100 200 300 400 500 600

20

30

40

50

R.A

. (%)E

long

atio

n (%

)

Temperature °C

40

50

60

70

80

90

100

As-received 1000 h 5000 h 10000 h

As-received 1000 h 5000 h 10000 h

U.T.S.

Y.S.

Elongation

R.A.

0 100 200 300 400 500 600200

300

400

500

600

700

RT

RT

Str

ess

(MP

a)

48

Figure 27: Impact properties of Eurofer'97.

-150 -100 -50 0 50 100

0

20

40

60

80

100

Brit

tle F

ract

ure

(%)

Temperature (°C)

-150 -100 -50 0 50 100-50

0

50

100

150

200

250

300

As-received Aged 500°C/5000h Aged 500°C/10000h Aged 600°C/5000h Aged 600°C/10000h

Ene

rgy

(J)

As-received Aged 500°C/5000h Aged 500°C/10000h Aged 600°C/5000h Aged 600°C/10000h

49

Figure 28: Fracture toughness results and associated Master Curve for as received Eurofer'97.

ASTM standard E 1921-02.

-180 -160 -140 -120 -100 -80 -60

50

100

150

200

250

300

350

400

450

EUROFER-97 1/2TCT As Received

KJC

1T (

MP

a√m

)

Temperature (ºC)

T0 = -129 ºC

MasterCurve

95%

5%

50

Figure 29: Fracture toughness results and associated Master Curve for Eurofer'97 aged at

600ºC/10000 hours. ASTM standard E 1921-02.

-180 -160 -140 -120 -100 -80 -60

50

100

150

200

250

300

350

400

450 EUROFER-97 1/2TCT Aged 600 °C, 10000 h

Temperature (ºC)

T0 = -122ºC

MasterCurve

95%

5%

K JC

1T M

Pa√

m

-180 -160 -140 -120 -100 -80 -60

50

100

150

200

250

300

350

400

450 EUROFER-97 1/2TCT Aged 600 °C, 10000 h

Temperature (ºC)

T0 = -122ºC

MasterCurve

95%

5%

EUROFER-97 1/2TCT Aged 600 °C, 10000 h

Temperature (ºC)

T0 = -122ºC

MasterCurve

95%

5%

K JC

1T M

Pa√

m

51

Figure 30: Strain- life data of Eurofer'97.

100 1000 10000

1

Test Temperature = 500°C

∆εt(%

)

Nf

Eurofer'97. As-received Eurofer'97. Aged 500°C/5000h

52

Figure 31: Relations between plastic and elastic strain and number of cycles to failure of

Eurofer'97 in the as-received condition.

1000 10000

0,1

1

∆εt

∆εe

∆εp

Eurofer'97 As-received. Test Tª = 500°C

∆εt=50.13N

f

(-0.62)+0.53Nf

(-0.06)

∆εt(%

)

Nf

53

Figure 32: Cyclic stress response curve of Eurofer'97 on as-received condition at different

strain ranges.

1 10 100 1000 10000

200

250

300

350

400

450

Pea

k st

ress

(M

Pa)

Number of cycles

∆εt=0.4%

∆εt=0.7%

∆εt=1%

∆εt=1.5%

54

Figure 33: Cyclic stress response curve of Eurofer'97 aged at 500°C/5000 h at different strain

ranges.

1 10 100 1000

200

250

300

350

400

450

Pea

k S

tres

s (M

Pa)

Number of cycles

∆εt = 0.7%

∆εt = 1%

∆εt = 1.5%

55

Figure 34: Typical creep and creep rate curves for the Eurofer'97 steel.

0 500 1000 1500 2000 2500 3000

1E-5

1E-4

1E-3

0 500 1000 1500 2000 2500 30000

5

10

15

20

0 200 400 600 800 1000 1200 14001E-5

1E-4

1E-3

0,01

0 200 400 600 800 1000 1200 14000

5

10

15

20

0 500 1000 1500 2000

1E-5

1E-4

1E-3

0,01

0,1

1

0 500 1000 1500 20000

5

10

15

20

25

450°C, 330MPATr = 2802 h

εmin = 1.22E-5 h-1

Stra

in (%

)C

reep

rat

e (h

-1)

500°C, 270MPATr = 1371 h

εmin = 2.69E-5 h-1

550°C, 190MPATr = 2196 h

εmin = 1.82E-5 h-1

0 1000 2000 3000 4000

1E-5

1E-4

1E-3

0 1000 2000 3000 40000

5

10

15

20

25

0 1000 2000 3000 4000

1E-5

1E-4

1E-3

0 1000 2000 3000 40000

5

10

15

20

εmin = 5.26E-6 h-1

600°C, 110MPATr = 4261 h

650°C, 60MPATr = 3968 h

εmin = 9.29E-6 h-1

Stra

in (%

)C

reep

rat

e (h

-1)

Time to rupture (h)

Time to rupture (h)

56

Figure 35: Variation of creep rupture strength against creep lifetime.

40

60

80

100

200

400

100 1000 10000

650°C

600°C

500°C

550°C

450°C

Time to rupture (h)

Str

ess

(MP

a)

57

Figure 36: Creep strength of Eurofer'97 steel.

40

60

80

100

200

400

23 24 25 26 27 28 29 30 31

200 MPa,500°C, Tr = 15470 h

Str

ess

(MP

a)

P = T(30+log t).10-3

Eurofer'97 As-received

58

Figure 37: Minimum creep rate against stress for the Eurofer'97 steel. Full lines are fitted by Norton equation.

1E-7

1E-6

1E-5

1E-4

1E-3

60 80 100 200 400

650°C

600°C 550°C

450°C

500°C

Stress (MPa)

Min

imum

cre

ep r

ate

(h-1)

59

Figure 38: Representative sub-grain structure of the creep specimens tested at short times. This image correspond to specimen C9S (650°C/70MPa/938 h).

60

Figure 39: Microstructural characterization in the creep sample tested at 500°C, 200 Pa and Tr = 15470 h. a) TEM micrograph and b) X-Ray

Diffraction pattern performed on the extracted residue showing M23C6 and M6C (Fe3W3C).

M23C6

M6C

M6C

M23C6

(a)

0 2 0 40 60 8 0 1000

500

1000

1500

2000

M6C

(F3W

3C)

M6C

(F 3

W3C

)

M6C

(F 3

W3C

)

Ta

CTa

C

M2

3C6

M23

C6

M23C

6

TaC

M6C

(F 3

W3C

)

M6C

(F3W

3C)

Cou

nts

2 T h e t a ( ° C )

(b)

61

Figure 40: EDS analyses (thin foil) of a) M6C and b) M23C6 precipitates.

Energy (keV)

Cou

nts

M6C

(a)

Cou

nts

Energy (keV)

M23C6

(b)

62

Figure 41: Phase extraction results of the Eurofer'97 and F-82H steels in the as-received

condition.

Eurofer'97 (2.5 %)* F-82H (1.5 %)*0

10

20

30

40

50

60

70

* % extracted residue

Con

cent

ratio

n (a

t %)

Material

Cr Fe W V Ta

63

Figure 42: Tensile properties of several ferritic/martensitic steels (Ref. 31).

Eurofer'97

Temperature °C

U.T

.S (M

Pa)

Eurofer'97Eurofer'97

Temperature °C

U.T

.S (M

Pa)

Temperature °C

Y.S

(MPa

)

Eurofer'97

Temperature °C

Y.S

(MPa

)

Eurofer'97Eurofer'97

Eurofer'97

Temperature °C

U.T

.S (M

Pa)

Eurofer'97Eurofer'97

Temperature °C

U.T

.S (M

Pa)

Temperature °C

Y.S

(MPa

)

Eurofer'97

Temperature °C

Y.S

(MPa

)

Eurofer'97Eurofer'97

64

Figure 43: Impact properties of several ferritic/martensitic steels (Ref. 31).

Eurofer'97Eurofer'97Eurofer'97

Eurofer'97Eurofer'97

Eurofer'97Eurofer'97Eurofer'97

Eurofer'97Eurofer'97

65

Figure 44: KJm for a plate of the Eurofer’97 as a function of B for tests at –142°C on

specimens with W = 14 mm (Ref. 53).

Figure 45: The F82H KJmr data versus temperature (Ref.56).