9-12% Cr heat resistant steels : alloy design, TEM characterisation ...

9-12% Cr heat resistant steels: alloy design, TEM characterisation

of microstructure evolution and creep response at 650°C

Dissertation

Zur Erlangung des Grades

Doktor-Ingenieur

der

Fakultät für Maschinenbau

der Ruhr-Universität Bochum

von

David Rojas Jara

aus Constitución, Chile

Bochum 2011

Dissertation eingereicht am: 11th January

Tag der mündlichen Prüfung: 21st March

Erster Referent: Prof. Dr. Anke Kaysser-Pyzalla

Zweiter Referent: Prof. Dr. Gerhard Sauthoff

Executive summary

This work was carried out aiming to design and characterise 9-12% Cr steels with tailor-

made microstructures for applications in fossil fuel fired power plants. The investigations

concentrated in the design and characterisation of heat resistant steels for applications in

high oxidising atmospheres (12% Cr) and 9% Cr alloys for components such as rotors

(P91).

ThermoCalc calculations showed to be a reliable tool for alloy development. The

modeling also provided valuable information for the adjustment of the processing

parameters (austenisation and tempering temperatures).

Two 12% Cr heat resistant steels with a fine dispersion of nano precipitates were

designed and produced supported by thermodynamic modeling (ThermoCalc). A detailed

characterisation of the microstructure evolution at different creep times (100 MPa /

650°C / 8000 h) was carried out by scanning transmission electron microscopy (STEM).

The results of the microstructure analysis were correlated with the mechanical properties

in order to investigate the influence of different precipitates (especially M23C6 carbides)

on the creep strength of the alloys. Precipitation of Laves phase and Z-phase was

observed after several hundred hours creep time. Very few Z-phase of the type

Cr(V,Ta)N nucleating from existing (V,Ta)(C,N) was observed. Both alloys show growth

and coarsening of Laves phase, meanwhile the MX carbonitrides present a very slow

growth and coarsening rate. Alloys containing Laves phase, MX and M23C6 precipitates

show best creep properties.

The influence of hot-deformation and tempering temperature on the microstructure

evolution on one of the designed 12% Cr alloys was studied during short-term creep at

80-250 MPa and 650°C. Quantitative determination of dislocation density and sub-grain

size in the initial microstructure and after creep was investigated by STEM combined

with the high-angle annular dark-field detector (HAADF). A correlation between

microstructure evolution and creep response was established. All crept samples showed a

significant increase of sub-grain size and a reduction of the dislocation density. Hot-

deformed samples showed better creep strength than non hot-deformed samples, due to

homogenisation of the microstructure. The tempering temperature affected the dislocation

density and the sub-grain sizes, influencing the creep behaviour.

9% Cr alloys were designed supported by ThermoCalc. Two sets of alloys were

produced: 9% Cr alloys with 0.1 % C and 0.05% C and 9% Cr alloys containing ~ 0.03%

Ti again with 0.1% C and 0.05% C (always wt%). Microstructure investigations showed

good agreement with the predicted phases of the thermodynamic modeling. The volume

fraction of precipitated M23C6 carbides is directly related to the carbon content of the

alloys. Hardening of the Ti-containing alloys by precipitation of fine dispersed Ti-based

MX particles was achieved. The precipitation of these carbides was limited to the

austenisation and tempering treatment used. The microstructure evolution (sub-grain and

particle size) during creep at 650°C / 100MPa was investigated by STEM-HAADF. The

sub-grain size evolution and the coarsening of precipitates (MX carbonitrides, M23C6 and

Laves phase) were more pronounced for Ti-containing alloys. 9Cr alloys without Ti and

with low carbon content presented the highest creep strength of all investigated alloys.

i

Table of contents 1. Introduction. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1

2. State of the art and objectives of the work. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3

2.1. Applications. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3 2.2. 9-12% Cr heat resistant steels. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4 2.3. Objectives. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6

3. Metallurgy of 9-12% Cr steels. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 10

3.1. Fundamentals of creep. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17 3.2. Microstructural changes during creep. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 20 3.3. Strengthening mechanisms. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 24

4. Methods. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 27

4.1. Thermodynamic modeling (ThermoCalc) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 27 4.2. Material preparation. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 33 4.3. Optical microscopy characterisation. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 33 4.4. Transmission electron microscopy characterisation. . . . . . . . . . . . . . . . . . . . . . . . . . 34

• Quantitative determination of precipitates. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 34 • Quantitative determination of dislocation density. . . . . . . . . . . . . . . . . . . . . . . 35 • Quantitative determination of sub-grain size. . . . . . . . . . . . . . . . . . . . . . . . . . . 38

4.5. Mechanical tests. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 38 • Creep tests. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 38 • Hardness. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 38

5. Results of 12% Cr steels. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 39

5.1 Alloy design and characterisation. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 40

5.1.1. Alloy design of 12% Cr steels supported by ThermoCalc. . . . . . . . . . . . . . . . . . 40 • Influence of Co and W on the microstructure formation. . . . . . . . . . . . . . . . . . 40 • 12Cr4CoWTa and 12Cr2CoWV design. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 43

5.1.2. Alloy production. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 47 5.1.3. Microstructure evolution (precipitates quantification) . . . . . . . . . . . . . . . . . . . . . 48

• Alloy 12Cr4CoWTa-780 HD. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 48 • Alloy 12Cr2CoWV-780 HD. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 51 • Creep results. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 54

5.2. Influence of processing parameters. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 55

5.2.1. Alloy processing. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 55 5.2.2. Initial microstructure after tempering (dislocation density and sub-grain size). . 56 5.2.3. Microstructure features at the initial state. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 58 5.2.4. Microstructure evolution during creep (dislocation density and sub-grain size). 59 5.2.5. Influence of hot deformation on creep strength. . . . . . . . . . . . . . . . . . . . . . . . . . 61 5.2.6. Influence of tempering temperature on creep strength. . . . . . . . . . . . . . . . . . . . . 62

ii

6. Discussion of 12% Cr steels. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 66

6.1. Alloy design and characterisation. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 66

6.1.1. Microstructure evolution. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 66 • Alloy 12Cr4CoWTa-780 HD. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 66 • Alloy 12Cr2CoWV-780 HD. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 67 • Creep. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 69

6.2. Influence of processing parameters. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 70 6.2.1. Initial microstructure after tempering (dislocation density and sub-grain size) . . 70 6.2.2. Influence of hot-deformation on initial microstructure. . . . . . . . . . . . . . . . . . . . . 71 6.2.3. Influence of tempering temperature on initial microstructure. . . . . . . . . . . . . . . . 72 6.2.4. Influence of hot deformation on creep strength. . . . . . . . . . . . . . . . . . . . . . . . . . . 73 6.2.5. Influence of tempering temperature on creep strength. . . . . . . . . . . . . . . . . . . . . . 74 6.2.6. Conclusions for the studied 12% Cr steels. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 75

7. Results of 9% Cr steels. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 78

7.1. Thermodynamic calculations of 9% Cr steels. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 78 • 9CrTi-H and 9CrTi-L design. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 79 • 9Cr-H and 9Cr-L design. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 81

7.2. Alloy production. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 84 7.3. Microstructure evolution (precipitates and sub-grain size). . . . . . . . . . . . . . . . . . . . . 86

• Alloys 9CrTi-H and 9CrTi-L (Influence of Ti). . . . . . . . . . . . . . . . . . . . . . . . . . 86 • Alloys 9Cr-H and 9Cr-L (Influence of C). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 92 • Creep results. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 97

8. Discussion of 9% Cr steels. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 99

8.1. Microstructure evolution (precipitates and sub-grain size) . . . . . . . . . . . . . . . . . . . . . 99 • Alloys 9CrTi-H and 9CrTi-L (Influence of Ti) . . . . . . . . . . . . . . . . . . . . . . . . . . 99 • Alloys 9Cr-H and 9Cr-L (Influence of C) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 102• Creep. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 104

8.2. Conclusions of studied 9% Cr steels. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 105

9. Final conclusion and perspectives. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 107

10. References. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 110

iii

List of figures Fig. 1-1: Net efficiency as function of steam pressure and temperature. . . . . . . . . . . . . . . . .

1

Fig. 2-1: Photographs of different applications of the 9-12% Cr steels in the fossil fired power plant industry [29]. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

3

Fig. 3-1: Phase diagram for Fe-xCr-0.1C calculated with ThermoCalc TCFe6 (γ = austenite, α = ferrite and σ = sigma phase). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

10

Fig. 3-2: Formation of a Z-phase particle by Cr diffusion from the ferritic matrix [19]. . . . .

17

Fig. 3-3: Schematic creep curve of engineering steel under constant tensile load and constant temperature. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

17

Fig. 3-4: Schematic illustration of microstructure of 9-12% Cr steel after tempering (Internal interfaces and precipitates) (modified from [101]). . . . . . . . . . . . . . . . . . . . . . . . . .

21

Fig. 3-5: Schematic illustration of microstructure evolution of 9-12% Cr steels after creep exposure (coarsening of internal interfaces and precipitation of more stables phases) (adapted from [101]). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

21

Fig. 4-1: Modification introduced by the regular term: Gibbs energy of mixing with its two contributions as a function of x. If T < w/2R, the curve has two points of inflection between which there is a miscibility gap; this is the case illustrated here. . . . . . . . . . . . . . . . . . . . . . .

30

Fig. 4-2: Simple body-centred cubic structure with preferential occupation of atoms in the body-centre and corner positions. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

32

Fig. 4-3: Scheme of bright field, annular dark field and high-angle annular dark field detectors of a STEM. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

35

Fig. 4-4: Contrast of dislocations for multi-beam case in different zone axes (sample 12Cr4CoWTa-780 NHD at initial stage). For the zone axes [131] (A) and [531] (A) a reduced contrast is obtained. For the multi-beam case of low index [110] (B) high contrast is achieved and dislocations are highlighted. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

37

Fig. 5-1: ThermoCalc phase fields as a function of the Co content for the reference alloy (F=ferrite and A= austenite, ThermoCalc TCFe6). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

42

Fig. 5-2: ThermoCalc phase fields as a function of the W content for the reference alloy (F=ferrite and A= austenite, ThermoCalc TCFe6). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

42

Fig. 5-3: 12% Cr steels heat treatment scheme. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

48

Fig. 5-4: TEM image of alloy 12Cr4CoWTa-780 HD in the initial stage. Black arrows show MX particles, white arrows show Laves Phase. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

49

Fig. 5-5: TEM image of alloy 12Cr4CoWTa-780 HD in the initial stage. C rich Ta-MX (white arrows) and N rich Ta-MX particles (black arrows). . . . . . . . . . . . . . . . . . . . . . . . . . .

49

iv

Fig. 5-6: TEM image of sample 12Cr4CoWTa-780 HD after 3,650 h under creep condition 650°C at 100 MPa. Laves phase (black arrows) and MX precipitates (white arrows) are present in the microstructure. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

50

Fig. 5-7: Diffraction pattern of Laves phase in sample 12Cr4CoWT-780 HD after 3,650 h creep. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

50

Fig. 5-8: EDS spectrum of Laves phase in sample 12Cr4CoWTa-780 HD after 3,650 h creep. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

50

Fig. 5-9: TEM image of alloy 12Cr2CoWV-780 HD in the initial stage. M23C6 carbides (white arrows) and MX precipitates (black arrows). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

51

Fig. 5-10: EDS spectrum of MX of the type (V,Ta)(C,N) showing V and Ta as main elements. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

51

Fig. 5-11: Diffraction pattern of M23C6 carbide in the initial stage of alloy 12Cr2CoWV-780 HD (fcc crystal structure). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

52

Fig. 5-12: EDS spectrum M23C6 precipitate in the initial stage of alloy 12Cr2CoWV-780 HD showing the main elements Cr, V, W and Fe. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

52

Fig. 5-13: TEM image of alloy 12Cr2CoWV-780 HD after 6,150 h creep. M23C6 carbide (white arrows) and Laves phase (black arrows) are observed. . . . . . . . . . . . . . . . . . . . . . . . .

53

Fig. 5-14: TEM image of alloy 12Cr2CoWV-780 HD after 6,150 h creep showing a nano-sized MX particle of the type (V,Ta)(C,N) (white arrow). . . . . . . . . . . . . . . . . . . . . . . . . . . .

53

Fig. 5-15: STEM image of alloy 12Cr2CoWV-780 HD after 6,150 h creep showing Laves phase and M23C6 carbides (black arrows) and Z-Phase (white arrow). The white points in the Z-phase indicate the EDS measurements. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

53

Fig. 5-16: EDS of Z-phase particle shown in Fig. 5-15. Three measurements were carried out in this phase (white points). The main elements are V, Cr, Ta and N. . . . . . . . . . . . . . . .

53

Fig. 5-17: Results of the tensile creep tests showing times to rupture as a function of applied stress for alloys 12Cr4CoWTa-780 HD and 12Cr2CoWV-780 HD at 650°C. An increased creep strength for alloy 12Cr2CoWV-780 HD can be observed. Results of creep tests of a P92 steel under similar conditions [41] are shown as reference. . . . . . . . . . . . . . . .

54

Fig. 5-18: Alloy 12Cr4CoWTa heat treatment scheme. . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

56

Fig. 5-19: STEM-HAADF image of the initial microstructure of sample 12Cr4CoWTa-780 HD (inverse contrast). The square area was amplified for better observation of the internal interfaces. A prior austenite grain boundary (dashed line), prior martensite laths (dotted lines) and sub-grain boundaries (full lines) are shown. . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

57

Fig. 5-20: Montage of STEM-HAADF images of the initial microstructure of sample 12Cr4CoWTa-780 HD (inverse contrast). Precipitates, sub-grains and dislocations are observed. For each picture of the montage, a multi-beam case with low index zone axis was adjusted in order to highlight the dislocation inside the sub-grains. . . . . . . . . . . . . . . . .

57

v

Fig. 5-21: STEM-HAADF image of sample 12Cr4CoWTa-680 HD after creep (2,875 h / 80MPa) showing the interaction of dislocations as well as with precipitates (A). Same image with inverse contrast for better observation of dislocation networks (B). . . . . . . . . . .

59

Fig. 5-22: Montage of STEM-HAADF images for (A) sample 12Cr4CoWTa-780 HD after creep (1,121 h / 145MPa / 650°C) and (B) sample 12Cr4CoWTa-780 HD after creep (3,650 h / 80MPa / 650°C). Smaller sub-grain sizes and higher dislocation density are observed for the sample with shorter creep time (A). White arrows indicate the size of some sub-grains for comparison. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

60

Fig. 5-23: Tensile creep test curves comparing the creep strength of sample 12Cr4CoWTa-780 NHD and sample 12Cr4CoWTa-780HD at 145 MPa / 650°C. (A) strain vs. time (B) creep rate vs. strain. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

62

Fig. 5-24: Tensile creep test curves comparing sample 12Cr4CoWTa-780 HD and sample 12Cr4CoWTa-680 HD at 80 MPa / 650°C. (A) Strain vs. time (B) creep rate vs. strain. . . .

63

Fig. 5-25: STEM-HAADF image of (A) sample 12Cr4CoWTa-680 HD after creep (2,875 h) and (B) sample 12Cr4CoWTa-780 HD after creep (3,650 h). Smaller sub-grain sizes are observed on sample (A). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

64

Fig. 5-26: Time to rupture as a function of applied tensile stress for 12Cr4CoWTa-680 HD and 12Cr4CoWTa-780 HD. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

65

Fig. 6-1: Formation process of M23(C,B)6 during heat treatment [38]. . . . . . . . . . . . . . . . . . .

70

Fig. 6-2: STEM-HAADF images of the initial microstructure of sample 12Cr4CoWTa-780 HD (a) and sample 12Cr4CoWTa-780 NHD (b). The microstructure of the hot-deformed sample shows a uniform distribution of precipitates compared to the non hot-deformed case. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

71

Fig. 7-1: ThermoCalc phase diagram for alloys 9CrTi-H and 9CrTi-L (F=ferrite and A= austenite, ThermoCalc TCFe6). The austenisation temperature and the tempering temperature are indicated in the phase diagram by black circles for each alloy. Ti-MX denotes the Ti-rich phase which contains N, C and few Nb, whereas Nb-MX are Nb-rich particles with C and N and also few amounts of Ti and Cr. . . . . . . . . . . . . . . . . . . . . . . . . . .

79

Fig. 7-2: ThermoCalc phase diagram for alloys 9Cr-H and 9Cr-L (F=ferrite and A= austenite, ThermoCalc TCFe6). The austenisation temperature and the tempering temperature are indicated in the phase diagram by black circles for each alloy. V-MX is V-rich phase containing Nb, N and C and few Fe and Cr, whereas Nb-MX denotes Nb-rich particles with C, Cr and N and also few amounts of V. . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

82

Fig. 7-3: 9% Cr steels heat treatments scheme. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

84

Fig. 7-4: STEM-HAADF micrographs of initial microstructure of alloy 9CrTi-H (A) and alloy 9CrTi-L (B). White arrows point at the M23C6 precipitates in both micrographs. Alloy 9CrTi-L shows large particles rich in W and Fe (possibly FeW2B). . . . . . . . . . . . . . . . . . . .

87

vi

Fig. 7-5: STEM-HAADF micrograph of the initial microstructure of alloy 9CrTi-H (A) (nano-sized Nb-MX precipitates are indicated by white arrows; black arrows point at the Ti-MX particles). EDS spectrum of the encircled Nb-MX precipitate (B). . . . . . . . . . . . . . .

88

Fig. 7-6: STEM-HAADF micrograph of initial microstructure of alloy 9CrTi-L (A) showing Ti-rich MX precipitates (white arrows), the M23C6 precipitates are pointed at by black arrows and EDS spectrum of the encircled Ti-MX particle (B). . . . . . . . . . . . . . . . . . .

88

Fig. 7-7: STEM-HAADF micrograph initial microstructure of alloy 9CrTi-L (A) showing a large Ti-rich precipitate and the M23C6 precipitates (white arrows) and EDS spectrum of the Ti-rich particle (B). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

89

Fig. 7-8: STEM-HAADF micrograph of alloy 9CrTi-H (A) after creep (7,253h / 101MPa / 650°C) and STEM-HAADF micrograph of alloy 9CrTi-L (B) after creep (2,154h / 101MPa / 650°C) with M23C6 precipitates (white arrows) and Laves phase (black arrows). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

90

Fig. 7-9: STEM-HAADF micrograph of alloy 9CrTi-L (A) after creep (2,154h / 101 MPa / 650°C), diffraction pattern of the encircled Laves phase particle (B), and EDS spectrum of Laves phase particle (C). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

91

Fig. 7-10: Tensile creep curves comparing the creep strength of alloys 9CrTi-H and 9CrTi-L at 101 MPa / 650°C, (A) strain vs. time, (B) creep vs. strain. . . . . . . . . . . . . . . . . . . . . . . .

92

Fig. 7-11: STEM-HAADF micrographs of initial microstructure of alloy 9Cr-H (A) and alloy 9Cr-L (B) with M23C6 precipitates (white arrows). . . . . . . . . . . . . . . . . . . . . . . . . . . . .

93

Fig. 7-12: STEM-HAADF micrograph of the initial microstructure of alloy 9Cr-L (A) with Nb-MX particles and EDS spectrum of the encircled Nb-MX particle (B). . . . . . . . . . . . . . .

94

Fig. 7-13: STEM-HAADF micrographs of the initial microstructure of alloy 9Cr-L (A) with V-MX particles (white arrows) and Nb-MX particles (black arrows) and EDS spectrum of the encircled V-MX particle (B). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

94

Fig. 7-14: STEM-HAADF micrograph of alloy 9Cr-H (A) after creep (7,987h / 101MPa / 650°C) and STEM-HAADF micrograph of alloy 9Cr-L (B) after creep (10,168h / 125MPa / 650°C) white arrows indicate M23C6 precipitates; black arrows indicate Laves phase. . . . .

96

Fig. 7-15: (A) STEM-HAADF micrograph of alloy 9Cr-L after creep (10,168h / 125 MPa / 650°C). (B) Diffraction pattern of the encircled M23C6 particle. (C) EDS spectrum of the encircled M23C6 particle. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

96

Fig. 7-16: STEM-HAADF micrograph of sample 9Cr-L after 10,168 h creep at 650°C / 125MPa (inversed contrast). Black arrows point at Laves phase particles, white arrows indicate the M23C6 carbides. Sub-grains and dislocations are often pinned by the M23C6 carbides. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

97

Fig. 7-17: Results of the tensile creep tests at 650°C showing time to rupture as a function of applied stress for the four investigated alloys. The alloy 9Cr-L as well as 9CrTi-H and 9Cr-H show the highest creep strength. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

98

vii

List of tables Tab. 2-1: Nominal chemical composition and creep rupture strength at 600°C of the historical development of 9-12% Cr steels. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

5

Tab. 3-1: Unit cell parameter of MX precipitates in 9-12% Cr steels. . . . . . . . . . . . . . . . . .

15

Tab. 5-1: Calculated composition (wt%) of precipitates for alloy 12Cr4CoWTa with ThermoCalc. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

45

Tab. 5-2: Calculated composition (wt%) of precipitates for alloy 12Cr2CoWV with ThermoCalc. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

45

Tab. 5-3: Volume fraction % of precipitates for alloy 12Cr4CoWTa calculated with ThermoCalc at tempering temperature (780°C) and at the creep testing temperature (650°C) (MX are of the type Nb(C,N)). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

46

Tab. 5-4: Volume fraction % of precipitates for alloy 12Cr2CoWV calculated with ThermoCalc at tempering temperature (780°C) and at the creep testing temperature (650°C) (MX are of the type Nb(C,N)). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

46

Tab. 5-5: Analysed chemical composition of the model alloys investigated. . . . . . . . . . . . .

47

Tab. 5-6: Mean size of precipitates in alloy 12Cr4CoWTa-780 HD (time in hours and size in nanometers). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

49

Tab. 5-7: Mean size of precipitates in alloy 12Cr2CoWV-780 HD (time in hours and size in nanometers). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

51

Tab. 5-8: Quantitative determination of PAGS, sub-grain size, dislocation density and hardness at the initial stage. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

58

Tab. 5-9 Effect of hot-deformation: Comparison of sub-grain size, dislocation density and hardness for samples 12Cr4CoWTa-780 NHD and 12Cr4CoWTa-780 HD after creep. . . .

62

Tab. 5-10 Effect of tempering temperature: Comparison of sub-grain size, dislocation density and hardness for samples 12Cr4CoWTa-680 HD and 12Cr4CoWTa-780 HD after creep. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

64

Tab. 7-1: Austenisation temperatures from ThermoCalc. . . . . . . . . . . . . . . . . . . . . . . . . . . .

80

Tab. 7-2: Volume fractions of precipitates calculated with ThermoCalc for alloys 9CrTi-H and 9CrTi-L at 780°C and 650°C. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

81

Tab. 7-3: Volume fractions of precipitates calculated with ThermoCalc for alloys 9Cr-H and 9Cr-L at 780°C and 650°C. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

83

Tab. 7-4: Analysed chemical composition of the produced alloys (wt%). . . . . . . . . . . . . . .

85

Tab. 7-5: Sub-grain size and hardness of alloys 9CrTi-H and 9CrTi-L at initial stage. . . . .

86

viii

Tab. 7-6: Average size of precipitates in alloys 9CrTi-H and 9CrTi-L at initial stage (time in hours and size in nanometers). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

87

Tab. 7-7: Sub-grain size and hardness of alloys 9CrTi-H and 9CrTi-L after creep at 650°C. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

89

Tab. 7-8: Average size of precipitates in alloys 9CrTi-H and 9CrTi-L (time in hours and size in nanometers) after creep (650°C / 101MPa). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

90

Tab. 7-9: Sub-grain size and hardness of alloys 9Cr-H and 9Cr-L at initial stage. . . . . . . . .

92

Tab. 7-10: Average size of precipitates in alloys 9Cr-H and 9Cr-L at the initial stage (time in hours and size in nanometers). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

93

Tab. 7-11: Sub-grain size and hardness in alloys 9Cr-H and 9Cr-L after creep at 650°C. . .

95

Tab.7-12: Average size of precipitates from alloys 9CrTi-H and 9CrTi-L (time in hours and size in nanometers) under creep condition (9Cr-H 650°C / 101MPa and 9Cr-L 650°C / 125MPa). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

95

Introduction

1

1. Introduction

Coal fired steam power plants produce ca. 40% of the world electricity. They are

expected to continue to do so in the next 30 years, due to the increasing global demand

for electricity. Coal fired power plants are intensive sources of global CO2 emissions.

Any improvement in efficiency of the best available coal power plants can thus have a

large effect on world environment conditions [1,2,3]. The thermal efficiency of steam

power plants is to a large extent controlled by the achievable temperature and pressure in

the steam cycle, which are limited by the properties of the available construction

materials [4,5]. Key materials for further improvements are the martensitic 9-12% Cr

steels used for thick section boiler components and turbines [6,7]. They offer the best

combination of high creep strength, high resistance against thermal fatigue, high steam

oxidation resistance, low cost and good manufacturability [8,9,10]. A doubling of creep

strength has been achieved for these materials over the last decades, leading to increases

in steam parameters from 180 bar / 530-540°C to 300 bar / 600°C, corresponding to

roughly 30% reduction in specific CO2 emissions [11].

Fig. 1-1: Net efficiency as function of steam pressure and temperature.

As shown in Fig. 1-1 there is a permanent driving force to increase the working

temperature in order to improve the thermal efficiency of fossil-fired power plants.

Introduction

2

The working temperature and the exposure at stress during service promote

microstructural degradation reducing creep strength [12].

Worldwide research on enhancing the creep strength of 9-12% Cr steels for working

operations at 650°C has revealed the importance of taking into account microstructural

evolution phenomena during creep, such as precipitation and coarsening of carbonitrides

and intermetallic compounds, as well as coarsening of sub-grains [13,14].

Many attempts have been made in the last 10-15 years in order to develop strong

martensitic 11-12% Cr steels, but all of these steels failed in long-term creep [15,16,17].

It now seems clear that the failure is mainly due to unexpected precipitation of coarse Z-

phase (Cr(V,Nb)N nitrides), which dissolve the finely dispersed V and Nb rich MX

nitrides, which are essential to creep strength [ 18 , 19 ]. Recent investigations have

demonstrated that the Z-phase precipitation is strongly accelerated by high Cr contents in

these steels [20,21,22].

Newly developed steels contain 9-10% Cr and small additions of V, Nb and N, which

cause strengthening by finely dispersed V and Nb rich MX nitrides. They show stable

long-term creep behaviour up to 100,000 h at 600°C [23,24].

The 9-10% Cr steels have limited steam oxidation resistance, and in order to increase the

steam temperature above 600°C a higher Cr content of 11-12% is mandatory for

oxidation protection [25].

The steel development thus seems to have reached a critical point, where new thinking is

needed. Promising concepts for strong 9% Cr steels based on optimised boron and

nitrogen contents are under development [ 26, 27]. Alternatively, other strengthening

phases than the MX nitrides have to be investigated, which are not sensitive to high Cr

contents. A doubling of the creep strength is needed compared to the presently best

available steels in order to be able to achieve the main objective of steam power plants

which work at 325 bar / 650°C.

State of the art

3

2. State of the art and objective of the work

2.1. Applications

There are many potential applications for the 9-12% Cr steels, but the single largest use is

in the power generation industry. Specifically, they have found use in superheaters and

reheaters tubing, boilers, main steam pipes, bolting and turbine blades and rotors in fossil

fuelled power plants [28]. Fig 2-1 shows different applications of the 9-12% Cr steel in

the fossil steam power plants industry.

Fig. 2-1: Photographs of different applications of the 9-12% Cr steels in the fossil fired power plant industry

[29].

In nuclear power plants 9-12% Cr steels are used for steam generators and for

superheaters in gas - and sodium- cooled nuclear reactors. They are also being considered

for first-wall applications in fusion reactor systems [30]. Their high resistance to thermal

fatigue also makes these steels suitable as structural materials for fusion reactors [33].

In the petrochemical and chemical processing industries 9-12% Cr steels are used in

systems for hydrogen desulphurisation and in steam systems at high pressures and high

temperatures [33].

State of the art

4

2.2. 9-12% Cr heat resistant steels

The development of 9-12% Cr martensitic/ferritic steels is strongly motivated from both

economic and environmental perspectives to improve the efficiency of fossil fuel fired

power plants [31,32].

The design and production of 9-12% Cr martensitic/ferritic steels began in 1912 [33]

when Krupp and Mannesmann produced a 12% Cr steel containing 2-5% Mo. This type

of steel was used for steam turbine blades. Over the past 50 years 9-12% Cr steels have

found increasing applications in thick sections components of steam power plants [5,9].

The 12%CrMoV steels introduced in the power plants in the mid 1960s were developed

for thin and thick walled power station components. Their creep strength is based on

solution hardening and on the precipitation of M23C6 carbides. These steels have been

applied successfully in power stations over several decades [34].

In the late 1970s the P91 steel (modified 9Cr-1Mo) was developed for manufacturing of

pipes and vessels for fast breeder reactors [1,34]. This steel has been widely used for

pipes and small forgings in all new Japanese and European power plants with steam

temperatures up to 600°C. The increase of creep strength comparing to the 12%CrMoV

steel is caused by the formation of thermal stable V and Nb rich precipitates. A lower Cr

content of about 9%, in the range where the steel microstructure consists of tempered

bainite or tempered martensite, also contributes to the higher creep strength [35,36].

Years later, a Japanese steel development programme of Nippon Steel led to the

development of the P92 steel (NF616). With the P92 a further increase in stress rupture

strengthening was obtained by addition of 0.003% B, 1.8% W and a reduction of Mo

content from 1 to 0.5% [37]. The addition of B ensures thermally stables M23(C,B)6

precipitates, whereas the higher W content leads to higher amount of precipitated Laves

phase [38,39]. In Tab. 2-1 an overview of the historical development of the 9-12% Cr

heat resistant steels from 1950 to 2005 is shown.

State of the art

5

Tab. 2-1: Nominal chemical composition and creep rupture strength at 600°C of the historical development of 9-12% Cr steels [6]

Chemical composition (wt%) Rupture strength

at 600°C (MPa) country

steel

C Cr Mo Ni W V Nb N B 104h 105h

Basic steels

Germany 1 X22CrMoV12-1 0.22 12.0 1.0 0.50 - 0.30 - - - 103 59

UK 2 H 46 0.16 11.5 0.65 0.70 - 0.30 0.30 0.05 - 118 62

France 3 54T5 0.19 11.0 0.80 0.40 - 0.20 0.45 0.05 - 144 64

Japan 4 TAF 0.18 10.5 1.5 0.05 - 0.20 0.15 0.01 0.035 216 (150)

USA 5 11%CrMoVNbN 0.18 10.5 1.0 0.70 - 0.20 0.08 0.06 - 165 (85)

Advanced steels

USA 6 P91 0.1 9.0 1.0 <0.4 - 0.22 0.08 0.05 - 124 94

Japan 7 HCM 12 0.1 12.0 1.0 - 1.0 0.25 0.05 0.03 - - 75

Japan TMK 2 0.14 10.5 0.5 0.5 1.8 0.17 0.05 0.04 - 185 90

Europe 9 X18CrMoVNB 91 0.18 9.5 1.5 0.05 - 0.25 0.05 0.01 0.01 170 122

Europe 10 X12CrMoWVNbN 0.12 10.3 1.0 0.8 0.80 0.18 0.05 0.06 - 165 90

E911 0.11 9.0 0.95 0.2 1.0 0.20 0.08 0.06 - 139 98

Japan 11 P92 0.07 9.0 0.50 0.06 1.8 0.20 0.05 0.06 0.003 153 113

Japan 12 P122 0.1 11.0 0.40 1.0Cu <0.40 2.0 0.22 0.06 0.06 0.003 156 101

Japan 13 HCM 2S 0.06 2.25 0.20 0.2 - 0.25 0.05 0.02 0.003 - 80

Germany 14 7CrMoTiB 0.07 2.40 1.0 1.0 - 0.25 - 0.01 0.004 0.07Ti - 60

Cost 522 15 FB2 0.13 9.0 1.5 0.15 - 0.20 0.05 0.02 0.0085 - 125

Cost 522 16 CB2 0.12 9.0 1.5 0.15 - 0.20 0.06 0.02 0.011 - 125

Objectives

6

2.3. Objectives

This work was carried out aiming to design and characterise 9-12% Cr steels with tailor-

made microstructures for applications in fossil fuel fired power plants. The engineering

design of the steels was carried out in order to achieve high creep strengths at 650°C and

100 MPa. The creep strength was compared to conventional commercial creep resistant

steels (X20, P91 and P92) used in power plants. The observations and conclusions of this

work should give new insights to materials engineers working in the design and

processing of 9-12% Cr steels.

The alloy design was supported by thermodynamic calculations. ThermoCalc has been

used for calculations of the phase equilibria and the evaluation of phase stabilities, so that

the influence of composition (addition of elements) and heat treatments on the 9% and

12% Cr multi-component alloys could be modeled.

The microstructure analyses was carried out by modern experimental methods like

scanning transmission electron microscope (STEM) combined with the high-angle

annular dark-field (HAADF), which provide much better dislocation and sub-grain

contrast for a quantitative study of the microstructure.

The investigations concentrated in heat resistant steels for applications in high oxidising

atmospheres (12% Cr) and 9% Cr alloys for components such as rotors (P91).

12% Cr steels are mandatory for oxidation protection of thin-walled components, such as

boiler tubes, where the steam oxidation resistance of the 9-10% Cr alloys is

disadvantageous, due to their high oxidation rates at 650°C.

The aim of the design and characterisation of the 12% Cr steels is to study the effect of

different combinations of precipitates in the microstructure evolution and creep response

at 650°C. Despite 12% Cr creep steels such as X20 have been extensively investigated at

temperatures of 550°C, the microstructure features of the designed alloys differ

considerably from the commercial alloys. Hence the study should provide information for

the design of tailored microstructures of 12% Cr heat resistant steels at 650°C.

Objectives

7

In particular, the investigations should clarify the creep behaviour of 12% Cr steels with a

microstructure consisting of Laves phase and MX carbonitrides without M23C6 and to

compare the creep rupture life of such alloys with similar 12% Cr alloys, whose

microstructure consists of Laves phase, MX carbonitrides and M23C6. Several researches

have reported the benefits of the MX carbonitrides and M23C6 carbides [8,9,10] as well as

the B contribution on the reduction of the coarsening rate of M23C6 carbides [23,24]. The

Laves phase is a principal precipitate on these alloys [40]. The study should also clarify if

the formation of fine dispersed Laves phase is achievable and the coarsening of the Laves

phase during early stages of creep.

Another objective of this work is to observe the microstructure evolution of the designed

alloys regarding phase formation and phase evolution and to correlate this information

with the creep test results. Because of the alloy composition (high Cr, Ta-content) it is

expected to observe the formation of Z-phase, which is responsible for the sudden

breakdown of creep strength of 12% Cr steels after ca. 8,000 h at 650°C [19].

The influence of the processing parameters was investigated to determine their effects on

creep strength of the alloys. The evolution dislocations density and sub-grain size was

investigated by scanning transmission electron microscopy to clarify the effect of the hot-

forging and tempering temperature on microstructure evolution and creep strength. A

quantitative investigation of the microstructure should clarify the main contribution of the

processing steps to the creep strength of the alloys. It is expected that the hot-forging

process might improve the homogenisation and dispersion of the precipitates, due to the

introduction of more nucleation sites for precipitates, leading to an even distribution of

the particles (precipitates) during their precipitation.

9% Cr alloys were selected because of their best creep performance at temperatures ~

600°C for the 9-12% Cr steel family. As reported in [20] the relatively low Cr content of

9% Cr alloys reduces the driving force for the precipitation of the Z-phase, which is

detrimental for the creep strength [19]. Although the oxidation resistance is reduced by

the lower Cr content, the steam oxidation resistance of the 9% Cr steels at 600 and 650°C

Objectives

8

is good enough for large dimensional component parts, such as turbine rotors, which

present a small surface to volume ratio.

One of the novelties of these alloys compared to previous works on 9Cr-3W-3Co-BN

steels [26, 41] is the reduction of the W and Co content in order to reduce the cost of the

alloy. The influence of the reduction of both alloying elements on the creep strength at

650°C was investigated.

The 9% Cr alloys were designed to study the strengthening potential of Ti-MX

precipitates as well as the effect of the carbon concentration on the formation of phases

and its effect on the creep response at 650°C. Few investigations on the effect of Ti

additions on 9% Cr steels have been reported. Abe [42] stated that Ti additions may form

Ti-MX particles with very slow coarsening rate, which stabilises the martensitic lath

structure. But no further studies have been published regarding this topic. Consequently

detailed STEM investigations on the effect of Ti-MX precipitates on microstructure

evolution and creep strength were carried out.

Regarding the effect of the carbon concentration on the formation of precipitates, Abe

[41] reported a low coarsening rate of MX and M23C6 with very low C additions

(0.002 %) in a 9Cr-3W-3Co-BN steel. In this work the effect of two different carbon

contents (0.05 and 0.1%) was investigated regarding the identification of the precipitates

and the interaction of the particles with the sub-grain boundaries, but also on the volume

fraction of Laves phase and M23C6 carbides.

The general objectives of the work were:

12% Cr:

• To design 12% Cr alloys (12 elements alloy) assisted by thermodynamic

modeling with the aim of investigating the effects of the alloying elements on the

austenite field formation, the formation of M23C6 carbides, Laves phase and MX

carbonitrides.

Objectives

9

• To compare the microstructure formation, evolution and creep response of 12%

Cr steels containing MX, Laves phase with 12% Cr steels containing MX, Laves

phase and M23C6 carbides.

• To carry out a quantitative determination of the microstructure evolution

regarding MX carbonitrides, M23C6 carbides and Laves phase at the initial state

(after tempering) and after creep.

• To investigate the influence of hot-deformation and tempering temperature on the

microstructure formation and microstructure evolution during creep by

quantitative determination of dislocation density and sub-grain coarsening.

• To correlate the microstructural features of the designed 12% Cr alloys with their

creep strength.

9% Cr:

• To design 9% Cr alloys (11 elements alloy) containing MX carbonitrides, M23C6

carbides and Laves phase at 650°C assisted by thermodynamic modeling.

• To investigate the potential of using Ti-containing precipitates for strengthening

the 9% Cr martensitic/ferritic alloys.

• To investigate the influence of C content on the formation of MX carbonitrides,

M23C6 carbides and Laves phase as well as the effect on creep strength.

• To carry out a quantitative determination of the microstructure evolution

regarding MX carbonitrides, M23C6 carbides and Laves phase at the initial state

(after tempering) and after creep.

• To correlate the microstructural features of the designed 9% Cr alloys with their

creep strength.

Metallurgy of 9-12% Cr steels

10

3. Metallurgy of 9-12% Cr steels

The standard heat treatment of 9-12% Cr steels consists of austenisation and tempering.

The austenisation is usually carried out at high temperatures above the Ac1 temperature

(inside the austenitic loop, see Fig. 3-1) [43] in order to dissolve most carbides and

nitrides and to obtain a fully austenitic microstructure [44]. If ferrite is still present after

austenisation it is usually designated as δ-ferrite. After air cooling to room temperature,

the microstructure should become fully martensitic, with a high dislocation density [45].

As the steel is hard and brittle at this point, it is necessary to soften it by tempering.

Fig. 3-1: Phase diagram for Fe-xCr-0.1C calculated with ThermoCalc TCFe6 (γ = austenite, α = ferrite and σ

= sigma phase).

Martensite is a slightly distorted tetragonal form of the ferritic bcc crystal structure.

When austenite (fcc) is rapidly cooled, a direct (diffusionless) transformation into

martensite takes place by a shear process. In order to achieve optimum strength, it is

important that most of the austenite transforms into martensite on cooling to room

temperature [46]. This depends on the temperature where austenite is fully transformed

into martensite (Mf temperature). Mf should be above room temperature otherwise

retained austenite is present after cooling. This can lower the strength and lead to

untempered brittle martensitic after tempering, because a lowering of the C content in the


11

matrix due to carbide precipitation may raise the Mf. Normally air-cooling of 9-12% Cr

steels is sufficient for martensitic transformation, because the high level of Cr retards

diffusion of C, thus preventing the formation of ferrite [45].

The following equation gives a rough estimate of the effect of alloying elements (wt%)

on Ms temperature [47]:

Ms = 550°C – 450C – 33Mn – 20Cr – 17Ni – 10W – 20V – 10Cu – 11 Nb – 11Si + 15Co

(3.1)

The only element that raises the Ms temperature is Co, which also is an austenite former,

making it important in heavily alloyed steels.

Tempering is done in order to recover ductility from the hard and brittle martensite. The

temperature for tempering is usually in the range of 680-780°C depending on the

properties required [44]. Tempering temperatures in the low end of the range are used for

components like turbine rotors, where high tensile strength is required. The high end of

the range is used for pressurised components like steam pipes, where high toughness is

necessary [31].

Effect of alloying elements

The composition of the alloy has a great bearing on the phases formed and the sequences

of precipitation. It is useful to examine the effect of each element on the alloy properties

even though the problem became extremely complicated due to element interaction. The

elements used to make up the composition of these steels are added for many reasons.

Most elements are added to stabilise phases which are beneficial to creep resistance or to

suppress phases which are detrimental. Some are added for long-term solid solution

strengthening or to improve the corrosion resistance of the alloy. The influence on

microstructure of the main alloying elements can be summarised as follows:


12

Chromium: This is a ferrite stabilising element and a carbide former. Large Cr additions

of 9-12% provide the necessary oxidation and corrosion resistance, as well as the

strengthening of the material by precipitation of carbides. Additions larger than 11% Cr

were found to markedly increase corrosion resistance at 650°C [48]. However, high Cr

concentrations promote the formation of δ-ferrite [49]. In addition, several authors have

demonstrated that high Cr contents (around 12%) increase the driving force for

precipitation of Z-phase, which decreases the creep strength [19,20,50].

Cobalt: This element is used in order to stabilise the austenite field [40]. Co was found to

remain in solid solution in 12% Cr steels, even with concentrations up to 10 wt% [51].

Co also raises the Ms and the Curie temperature [40] and it is expected to slow down

diffusion processes, reducing the coarsening of the precipitates [52].

Copper: This is a very effective element to avoid the formation of δ-ferrite, which has a

detrimental effect on the mechanical properties of the steel [33]. At concentrations higher

than 0.5% Cu prevents a further sharp decrease of the Ac1 temperature [40]. Cu has a low

solubility in ferrite and forms Cu-rich precipitates, which may provide nucleation sites for

Laves phase formation [53].

Manganese: Mn stabilises the austenite but it is often found to have an adverse effect on

the creep strength of the creep resistant steels [54]. It has been found that increasing the

Mn content may increase the growth rate of M6C, an undesirable and coarse phase which

can remove W from solid solution and cause the dissolution of other important

precipitates such as M23C6 carbides and Laves phase [55].

Carbon: C occupies interstitial sites in both austenite and ferrite, with a greater solubility

in austenite. C stabilises the austenite relative to ferrite. It is essential for the formation of

carbides which causes the secondary hardening of the 9-12% Cr steels [41].

Nitrogen: N also occupies interstitial sites in the iron lattice and is an austenite stabiliser.

Increasing N stabilises MX precipitates, which are fine and generally desirable for creep


13

strength [56]. The addition of N to B containing steels should be restricted due to the

formation of BN, which can offset the beneficial effects of both B and N [57,58].

Silicon: Si is a ferrite stabilising element and can influence the kinetics of carbide

precipitation [ 59]. Si additions have been found to accelerate the precipitation and

coarsening of Laves phase [60,61]. Si additions can also promote the formation of δ-

ferrite phase so that austenite stabilising elements are often added purely to counteract

this effect [62]. Si is very important in the formation of protective oxidation layers [60].

Vanadium: V is a ferrite stabilising element and a strong carbide former [63]. It may also

combine with C and N to form fine V-rich carbides and carbonitrides (MX precipitates)

which significantly improve the long term creep strength [64].

Niobium: Nb is a ferrite stabilising element which forms stables MX precipitates with C

and N [65]. The effect of Nb depends on the austenisation temperature which governs the

amount of MX precipitates which is taken into solution [55,66].

Tantalum: Ta, as Nb, is a ferrite stabilising element which forms stable MX precipitates

with C and N [40]. The Ta-rich MX precipitates were found to be beneficial to the creep

rupture strength [67]. They are extremely stable and show slow coarsening rates during

creep.

Titanium: Ti is a strong nitride and carbide former and can improve the creep rupture

strength of ferritic steels [42]. Ti precipitates showed a very slow coarsening rate [68].

However, Ti in combination with N may promote the formation of large TiN, which may

decrease the creep strength [56].

Molybdenum: Mo is a ferrite stabilising element and also forms carbides [69]. Additions

of Mo can stabilise the M2X phase and the M23C6 phase [70]. Large additions (> 1%)

have been found to promote formation of M6C, Laves phase and δ-ferrite in 9% Cr steels

[61].


14

Tungsten: W is a strong carbide former which promotes the formation of Laves phase and

stabilises the ferrite [40]. W is also well known to increase the high temperature strength

via solid solution hardening [71,72]. W additions have been found to generally improve

the creep strength of ferritic heat resistant steels [41]. Abe et al. [73] reported that

increased W concentrations in 9% Cr alloys reduce the coarsening rate of M23C6

precipitates. This influences the coarsening of martensite laths due to the pinning effect

of M23C6, thus delaying recovery of the lath martensitic microstructure [74].

Boron: B stabilises the lath martensitic microstructure, due to the fact that B reduces the

Ostwald ripening rate of fine M23C6 carbides by the enrichment of B in the vicinity of

prior austenite boundaries [75,76].

Precipitates

Many of the precipitates which are formed in 9-12% Cr steels are metastable, and will

disappear with time. This fact is extremely relevant for the creep properties of 9-12% Cr

steels since the microstructure evolves during working condition at high temperatures.

Some of them have very short lifespan, which transform in the tempering process and are

replaced by more stable precipitates. The thermodynamic driving force for the metastable

precipitates is lower than for the stable precipitates. Their occurrence is kinetically

favoured either because the elements necessary for their formation are more readily

available in the matrix or their surface and/or strain energy terms involved in their

development are lower, e.g. nucleation is easier.

An example of a typical precipitation sequence in 9-12% Cr steel is as follows [47]:

M3C → M7C3 + M2X →M23C6 + MX → M23C6 + Z-phase

Here, there is a great difference in lifetime, as the early carbides may only exist for few

hours, while it can take decades for Z-phase to develop and to dissolve MX. Depending

on composition MX may be present during the entire life of certain steels.


15

In the following the main precipitates found in 9-12% Cr creep resistant steels are

described:

M23C6: M23C6 is a Cr-rich carbide which may also contain W, Mo, V, Fe and B [77,78].

The M23C6 has a cubic crystal structure (fcc space group Fm3m) with the lattice

parameter varying between 1.057 and 1.068 nm. M23C6 carbides are found in the early

stage of tempering, because they nucleate easily on the prior austenite grain boundaries

and martensite laths or block boundaries. After tempering the average size of the carbides

is about 100 nm [79], but the coarsening rate is comparatively high, decreasing their

influence on creep strength with time [13]. In B containing steels, B will dissolve in

M23C6 carbides and substitute for carbon, although only in very small quantities. The

enrichment of B in M23C6 carbide promotes the formation of intergranular M23(C,B)6

which may decrease the coarsening rate of the precipitate [39].

MX: The formation of MX precipitates occurs when strong carbides and/or nitrides

formers are added to the alloy (e.g. V, Nb, Ta, Ti) [64]. MX carbonitrides often have a

cubic NaCl-type structure. The lattice parameters of some of MX carbonitrides are shown

in Tab. 3-1. Often the lattice parameters have intermediate values, indicating the

existence of a solid solution between the different carbonitrides [80,81]. MX particles

usually form during tempering on dislocations within the matrix or on sub-grain

boundaries. They increase creep strength by pinning free dislocations and sub-grain

boundaries [82].

Tab. 3-1: Unit cell parameter of MX precipitates in 9-12% Cr steels.

Precipitates NbN NbC TiN TiC VC VN

a (nm) 0.439 0.447 0.424 0.433 0.417 0.413

Laves Phase: This is an intermetallic phase of the type (Fe,Cr)2(W,Mo) which may

precipitate in particular in Mo or W containing steels [41,61]. Laves phase also contains

minor amounts of Si. A Laves phase with a hexagonal crystal structure (space group P63)

with lattice parameters a=0.473 and c=9.772 nm [83] is usually found in 9-12% Cr steels.

Often the Laves phase does not nucleate during tempering as it is not stable at high


16

temperatures. The nucleation and growth rate is slow [ 84 , 85 ]. It precipitates

intergranularly during service exposure. During the growth step it becomes larger than

most other particles, but the coarsening rate is slower than M23C6 [13]. W-containing

Laves phase usually nucleates faster, thus becoming smaller and more finely distributed

as compared to the Mo-containing Laves phase [86]. The Laves phase has in the past

been blamed for the breakdowns in creep strength of several 9-12% Cr steels, arguing

that it removes W from the matrix, which would imply a loss of solid solution

strengthening by W. However, this explanation seems unlikely, as the precipitation

strengthening contribution of Laves phase should largely compensate the W depletion

[74].

Z-phase: Z-phase is probably the most stable nitride in 9-12% Cr steels during long-term

exposure in the temperature range 600-700°C [20]. It has an empirical formula of CrXN,

were X can be Nb, V or Ta. Jack et al. [87] first identified it as a tetragonal CrNbN

nitride. Danielsen et al. [88] identified a V containing modified Z-phase with a cubic

NaCl-type structure. In a recent work [89] Danielsen reported different Z-phase crystal

structures from several authors. According to Danielsen [90] the cubic structure was

found to coexist whit the tetragonal structure in the Z-phase. Further investigation [91]

showed that the cubic structure of Z-phase was predominant in samples which had been

exposed for relatively short times, while the tetragonal diffraction patterns became clearer

with longer exposure times.

The Cr content in 9-12% Cr steels has a strong influence on the precipitation kinetics of

Z-phase. 11-12% Cr steels have a much higher rate of Z-phase precipitation than 9% Cr

steels [92].

Z-phase precipitation causes the dissolution of MX carbonitrides (see Fig. 3-2), which are

beneficial to creep strength [19]. Hence progressive Z-phase precipitation causes

breakdown in creep strength [18]. High Cr steels show creep strength reductions

concurrent with Z-phase precipitation, whereas the steels with ~ 9% Cr and limited Z-

phase precipitation do not show effects. Therefore Z-phase directly contributes to a

reduction in the creep rupture strength on 11-12% Cr steels [93].


17

Fig. 3-2: Formation of a Z-phase particle by Cr diffusion from the ferritic matrix [19]

3.1. Fundamentals of Creep

Creep of materials is classically associated with time-dependent plastic flow under a

fixed stress at an elevated temperature, often greater than roughly 0.5 Tm, where Tm is

the absolute melting temperature [94]. Creep tests can be conducted either at constant

load or at constant stress. For experimental convenience, creep tests of engineering steels

are frequently conducted at constant tensile load and at constant temperature. The test

results can be plotted as creep curves, which represent graphically the time dependence of

strain measured over a reference or gauge length [95].

Creep of metals and alloys is generally described as a three stages phenomenon,

consisting of primary or transient creep, secondary or steady-state creep and tertiary or

accelerated creep. Fig. 3-3 shows a schematic diagram of a typical creep curve indicating

the three regimens.

Fig. 3-3: Schematic creep curve of engineering steel under constant tensile load and constant temperature.


18

In the primary creep or transient stage, the creep rate ε& (see equation 3.2) decreases with

time. The decreasing creep rate has been attributed to strain hardening or to a decrease in

free or mobile dislocation.

dtdεε =& (3.2)

In the secondary creep stage, the creep rate remains constant. This creep rate is

designated as a steady-state creep rate and is commonly attributed to a state of balance

between the rate of generation of dislocations contributing to hardening and the rate of

recovery contributing to softening. At high homologous temperatures, creep mainly

involves diffusion and hence the recovery rate is high enough to balance the strain

hardening and results in the appearance of secondary or steady-state creep. In the tertiary

creep stage, the creep rate increases with time until rupture at rupture time. It should be

considered that under a constant tensile load, the stress continuously increases as creep

proceeds because the cross-section decreases. For this reason a pronounced effect of

increase in stress on the creep rate appears in the tertiary creep stage. Necking of the

specimens before rupture causes a significant increase in stress. The increase in creep rate

with time in the tertiary creep stage is a consequence of increasing stress or of

microstructure evolution including damage. Microstructure evolution usually consists of

dynamic recovery, dynamic recrystallisation, coarsening of precipitates and other

phenomena, which cause softening and result in a decrease in resistance to creep.

Damage includes the development of creep voids and cracks, often along grain

boundaries [95,96].

Fig. 3-3 shows the idealised creep curve, however, engineering creep-resistant steels

sometimes exhibit complicated behaviour especially under low stress and long time

conditions, reflecting complex microstructural evolution during creep.

Under certain conditions, the secondary or steady-state creep stage may be absent; so that

the tertiary creep stage begins immediately after the primary creep stage. In this case the


19

minimum creep rate, ε& min, can be defined instead of the steady-stage creep rate. Similar

to the steady-stage creep rate, the minimum creep rate can be explained by the process

where hardening in the primary stage is balanced by softening in the tertiary stage. In

many cases, there is substantially no steady-state stage in engineering creep-resistant

steels and alloys. Many researchers have associated this phenomenon with an ever-

evolving microstructure during creep. This suggests that in engineering creep-resistant

steels there is no dynamic microstructural equilibrium during creep, which characterises

steady-state creep of simple metals and alloys. Therefore, the term “minimum creep rate”

has been favoured by engineers and researchers who are concerned with engineering

creep-resistant steels and alloys [97]. The stress dependence of minimum or steady-state

creep rate is usually expressed by a power law as:

nAσε =min& (3.3)

where n is the stress exponent.

The value A is determined as follows:

)/exp( RTQKA c−= (3.4)

where Qc the activation energy for creep, R the gas constant and the T, the absolute

temperature. The parameter K includes microstructure parameters such as grain size.

Equation (3.3) is often referred to as Norton’s law. It is well known that the minimum

creep rate is inversely proportional to the time to rupture tr as follows:

m

rcn tCRTQK )/()/exp(min =−= σε& (3.5)

where C is a constant depending on the total elongation during creep and m is a constant

often nearly equal to 1. Equation (3.5) is referred to as the Monkman-Grant relationship

which suggests that the minimum creep rate and the time to rupture are closely related.


20

3.2. Microstructural changes during creep

The martensitic/ferritic matrix of 9-12% Cr steels contains a high dislocation density and

various internal interfaces, such as prior austenite grain boundaries, prior packet or block

boundaries, prior martensite lath boundaries and sub-grain boundaries. Depending on the

alloy composition different kinds of precipitates are present in the microstructure after

tempering, e.g. M23C6 carbides and MX carbonitrides (see Fig. 3-5) [98,99]. The working

temperature (up to 650°C) and the exposure at stress during service promote

microstructural changes in the sub-grains and precipitates reducing creep strength

[100,101]. Coarsening of precipitates (e.g. Laves phase, M23C6) and precipitation of

undesirable phases (e.g. Z-phase) decrease the strength of the steels during creep (Fig. 3-

6).

The particle coarsening process follows the Ostwald ripening mechanism [31]. During

ripening, the average precipitate particle volume increases with time at elevated

temperatures. As a result, the spacing of the precipitates - in particular on dislocations -

increases and particle hardening decreases [102].

According to Blum et al. [103] the recovery of free dislocations may control the creep

rate during the primary creep, but that control probably shifts during creep to sub-grain

boundary processes. In particular, the migration of sub-grain boundaries may take care of

recovery of dislocations.

In Fig 3-4 and 3-5 a schematic illustration of the coarsening of the sub-grain boundaries

and the precipitation of more stables phases is shown [104].


21

Fig. 3-4: Schematic illustration of microstructure of 9-12% Cr steel after tempering (Internal interfaces and

precipitates) (adapted from [104]).

Fig. 3-5: Schematic illustration of microstructure evolution of 9-12% Cr steels after creep exposure

(coarsening of internal interfaces and precipitation of more stables phases ) (adapted from [104]).

Carbide coarsening (Ostwald ripening)

The coarsening process has a profound effect on the properties of materials. For example,

the coarsening process can control the size of precipitates in solid solution [105]. This

average particle size then sets the mechanical properties of the precipitation hardening


22

[72]. Coarsening typically occurs under conditions where the volume fractions of the

phases are nearly at their equilibrium values.

The driving force for the Ostwald ripening process is the reduction of the interface free-

energy of the material [106]. Since smaller particles in solution have a higher surface to

volume ratio than larger particles, smaller particles are less stable than larger particles of

the same material. An increase in the mean particle size will thus reduce the total free

energy of the system and this reduction in free energy is the driving force for the

coarsening reaction. The analysis of Lifshitz, Slyozov and Wagner [107] shows that the

average particle size increases with time as t1/3 and that the number of particle per unit

volume decays as t-1, while the volume fraction remains constant.

The Ostwald ripening of MaCb carbide in Fe-M-C alloys is described by the following

expression [108]:

tkrr 33

03 =− (3.6)

23 )(9/)(8 MpMMM uuaRTuVDbak −+= σ (3.7)

where r and r0 are the average particle size at the times t and t = 0, respectively, σ is the

interfacial energy of the carbide, V is the molar volume of the carbides, DM is the volume-

diffusion of metal M, uM and upM are the concentrations of M in the matrix and carbides,

respectively, R is the gas constant and T is the temperature [27].

Sub-grain coarsening

The tempering of the 9-12% Cr steels leads to precipitation of solute atoms and to

recovery of the dislocation cell structure [109,110] resulting in a sub-grain structure,

characterised by the frequency distributions of misorientations and of boundary spacings.

The sub-grains are bounded by the boundaries of the prior austenite grains, of blocks of

martensite laths of similar orientation, and of martensite laths and of sub-grains within


23

the laths [111]. The prior austenite grain and block boundaries are of the high-angle type

with misorientations across the boundaries generally lying above 10-15° [ 112]. The

martensite laths and sub-grains within the laths are of the low-angle type with lower

misorientations. In contrast to low-angle boundaries constituting planar dislocation

networks, high-angle boundaries interrupt the coherency of the lattices of the

neighbouring crystallites. They can not in general be penetrated by dislocations. High-

angle boundaries allow grains to slide relative to each other and provide a particularly

effective short circuit path for diffusion of atoms.

The sub-grain size w (mean linear intercept) approaches a steady state value in the course

of deformation which scales in inverse proportion to stress:

w∞ = kwb G/σ (3.8)

Here b is the length of the Burgers vector, G is the elastic shear modulus and kw is a

numerical factor which for steels was reported to lie at about 10. A typical value of the

initial sub-grain size w0 in tempered martensite is 0.4 μm [101,109].

Sub-grain coarsening occurs in creep when the initial sub-grain size is smaller than the

steady state sub-grain size according to Equation (3.8). This is the case in most

applications of tempered martensite steels.

Reduction of creep strength by sub-grain coarsening was confirmed in tests where sub-

grains were made to coarsen within a relatively short time, where changes in the

precipitate structure are negligible [100]. The coarsening was achieved by strain-

controlled cyclic straining at elevated temperature. When the maximum stress acting in

cyclic deformation is sufficiently low, the sub-grains grow fast with accumulating

inelastic strain towards the instantaneous stress-dependent value of the steady state sub-

grain size (Equation 3.8). The sub-grain coarsening caused the minimum creep rate to

increase by about one order of magnitude [100].


24

3.3. Strengthening mechanisms

The basic methods by which creep-resistant steels can be strengthened are solute

hardening, precipitation or dispersion hardening, dislocation hardening and boundary or

sub-boundary hardening [76, 113 ]. The creep strength is a combination of several

strengthening mechanisms so that it is often difficult to determine the single contribution

of each mechanism to the overall creep strength.

Solid solution hardening

For this strengthening mechanism, solute atoms of one element are added to another,

resulting in either substitutional or interstitial point defects in the crystal. The solute

atoms cause lattice distortions that impede dislocation motion, increasing the yield stress

of the material. Solute atoms have stress fields around them which can interact with those

of dislocations. The presence of solute atoms impart compressive or tensile stresses to the

lattice, depending on solute size, which interfere with nearby dislocations, causing the

solute atoms to act as potential barriers to dislocation propagation and/or multiplication

[114].

Substitutional solute atoms such as Mo and W, which have much larger atomic sizes than

those of the iron matrix, are effective solid solution strengtheners for both ferritic and

austenitic creep-resistant steels.

It should be noted that the contribution of solid solution hardening by Mo and W to the

overall creep strength of engineering creep-resistant steels is practically superimposed on

other strengthening mechanisms, for example precipitation hardening [73,113].

Precipitation or dispersion hardening

Creep resistant steels usually contain several types of carbonitrides (e.g. M23C6, MX) and

intermetallic compounds (Laves phase) in the matrix and at grain boundaries [115]. The


25

dispersion of fine precipitates stabilises free dislocations and the sub-grain structure

against recovery, which further enhances dislocation hardening and sub-boundary

hardening, respectively [64].

Several mechanisms have been proposed for the threshold stress, corresponding to the

stress required for a dislocation to pass through precipitate particles, such as the Orowan

mechanism [6]. The Orowan stress orσ is given as follows:

λσ /8.0 MGbor = (3.9)

where M is the Taylor factor (= 3), G is the shear modulus, b is the magnitude of the

Burgers vector and λ is the mean interparticle spacing [113].

The coarsening of fine precipitates and for some compositions the dissolution of MX to

form massive precipitates of Z-phase cause an increase in λ in Equation (3.9) and hence a

decrease in Orowan stress over long periods of time [76,113]. The coarsening and

dissolution of fine precipitates takes sometimes place preferentially in the vicinity of

grain boundaries during creep, which promotes the formation of localised weak zones and

promotes localised creep deformation near grain boundaries [24,116]. This results in

premature creep rupture.

Dislocation hardening

Dislocation hardening is an important means of strengthening steels at ambient

temperature. Dislocation hardening (σor) is given by:

2/1)(5.0 for MGb ρσ = (3.10)


26

where ρf is the free dislocation density in the matrix. Tempered martensitic 9-12% Cr

steels usually contain a high density of dislocations in the matrix even after tempering,

usually in the range of 1-10 × 1014 m-2 [117].

At elevated temperatures, cold working enhances softening by promoting the recovery of

excess dislocations and the recrystallisation of the deformed microstructure, causing a

loss of creep strength [118]. Dislocation hardening is useful for creep strengthening in the

short term but it is not useful for increasing long-term creep strength at elevated

temperatures.

Sub-boundary hardening

Tempered martensitic high Cr steels subjected to normalising and tempering are usually

observed to have a lath martensitic microstructure consisting of lath and block with a

high density of dislocations and a dispersion of fine carbonitrides along the lath and block

boundaries and in the matrix. The lath and block can be regarded as elongated sub-grains.

The lath and block boundaries provide the sub-boundary hardening given by:

sgsg Gb λσ /10= (3.11)

where λsg is the short width of elongated sub-grains [113].

The coarsening of the laths and blocks with creep strain, which mainly takes place in the

tertiary or acceleration creep region [119] and causes an increase in λsg of equation (3.11),

indicates the mobile nature of lath and block boundaries under stress. In the acceleration

creep of tempered martensitic 9% Cr steels, the progressive local coalescence of two

adjacent lath boundaries near the Y-junction causes the movement of the Y-junction,

resulting in the coarsening of the laths [41]. It is well known that polygon and sub-grain

boundaries free from precipitates in pure metals and solid solution alloys are highly

mobile under applied stress [120]. The movement of lath and block boundaries can

absorb or scavenge excess dislocations from inside the laths and blocks. This corresponds

to a dynamic recovery process, resulting in softening.

Methods

27

4. Methods

4.1. Thermodynamic modeling (ThermoCalc)

The alloy design was carried out based on metallurgy principles and assisted by

computational thermodynamics. The ThermoCalc software based on the CALPHAD

method has been employed for the design of the alloys [121].

ThermoCalc is a powerful software for thermodynamic calculations in multicomponent

systems. It is widely used for calculations of:

• Phase diagrams

• Thermochemical data such as enthalpies, heat capacity, and activities

• Solidification simulations with the Scheil-Gulliver model

• Assessment of experimental data

CALPHAD modeling

The use of phase diagrams has, for long, been seen as being rather academic, because

most real materials are multicomponent in nature, while phase diagrams are generally

used to represent binary or ternary systems.

The CALPHAD (Calculation of Phase Diagram) method has altered this point of view

because it is now possible to predict the phase behaviour of complex, multicomponent

systems, based on the extrapolation of thermodynamic properties. At the heart of this

method is the calculation of the Gibbs energy of a phase as a function of its composition

temperature and pressure. Within this approach, the problem of predicting equilibrium is

essentially mathematical, although far from simple due to the number of variables

involved in the minimisation process.

Methods

28

The models used for the Gibbs energy description vary with the nature of the phase

considered. In the following a brief introduction to the CALPHAD description of pure

substances, solutions and sublattice phases (which are the most commonly used in the

field of metallurgy) is given. The phases studied in the present work fall into these three

categories: MX precipitates (TiN, NbN, VN, etc) are modeled as pure substances,

complex carbides (e.g. M23C6) and ferrite are sublattice phases, while the liquid phase is a

random substitutional solution.

Pure substances

For a stoichiometric compound, it is sufficient to know the heat capacity together with

reference values to obtain the Gibbs energy value at any temperature. The SGTE

(Scientific Group Thermodata Europe) [122] databases store the coefficients for the heat

capacity at constant pressure Cp of numerous substances, written as a polynomial of

temperature (Eq. 4.1) together with values for ΔHf, the enthalpy of formation of the

substance, and S298 the entropy at 298 K. The coefficients are valid only within a given

range of temperature and the database provides parameters as a function of temperature

interval.

22)( −+++= DTCTBTATC p (4.1)

Random substitutional solutions

In random substitutional solutions, such as gases, or simple metallic liquids and solid

solutions, the components can mix on any spatial position available to the phase. The

Gibbs energy of a solution is traditionally decomposed according to:

xsmix

idealmix GGGG ++= 0

(4.2)

Methods

29

where 0G is the contribution to the Gibbs energy of the pure components, idealmixG the ideal

mixing contribution and xsmixG is the deviation from the ideal solution, also known as

excess Gibbs energy of mixing.

Ideal solutions

This is the simplest possible case. The interactions between the different elements are

identical and there is no enthalpy change when the solution is formed. The only

contribution to the Gibbs energy change is due to the increase of configurational entropy.

This term can be simply calculated using Stirling’s approximation for large factorials:

∑−=Δi

iiramdom xxRS ln (4.3)

where ix is the atomic fraction of the component i and R is the gas constant. The Gibbs

energy per mole of solution is therefore:

ii

iii

i xxRTGxG ln0 ∑∑ += (4.4)

Regular and non regular solutions

In most cases, there are interactions between the components of a phase. In the case of a

binary system AB, the regular model assumes that the total energy of solution can be

written as:

ABABAAAABBAAAA NNNNE εεε ++= (4.5)

where NAA is the number of AA pairs and εAA their bond energy. It can be demonstrated

that the enthalpy of mixing is:

Methods

30

)2)(1(2 BBAAABmix xxNzH εεε −−−=Δ (4.6)

where N is the number of atoms in solution, and z the coordination number of the

structure, that is the number of nearest neighbours of any atom.

For a regular solution, it is also assumed that the entropy of mixing is given by equation

(4.3) so that it does not contribute to the excess Gibbs energy of mixing, therefore:

)1( xwxG xsmix −= (4.7)

where )2(2 BBAAAB

zNw εεε −−= is a temperature dependent parameter on which

depends the behaviour of the solution (Fig. 4-1). Generalising this equation to a

multicomponent random solution the Gibbs energy per mole of substance can be written

as:

∑∑∑∑≥

++=i iJ

iJJiii

iii

i wxxxxRTGxG ln0 (4.8)

Fig. 4-1: Modification introduced by the regular term: Gibbs energy of mixing with its two contributions as a

function of x. If T < w/2R, the curve has two points of inflection between which there is a miscibility gap; this

is the case illustrated here.

However, the later assumes that interactions are composition independent, which is not

realistic in most cases. The sub-regular model, proposed by Kaufman and Berstein

Methods

31

introduces a linear composition dependency and expresses the excess Gibbs energy of

mixing as:

∑∑ +=≥i

jiiJ

iJi

iiJJi

xsmix xwxwxxG )( (4.9)

This is generalized to any composition dependency in the Redlich-Kister power series

which expresses xsmixG as:

∑∑ ∑ +=≥i

vj

iJi

v

viJJi

xsmix xxwxxG )( (4.10)

For example in SGTE databases the individual parameter w is written as:

w = A+BT+CT lnT+DT2 (4.11)

and these coefficients are stored for each viJw

Sublattice models

The different expressions for xsmixG presented above are all for solutions where the

components can mix freely on the sites available for the phase (see Fig. 4-2). In many

cases, however different components mix on different sublattices, as with ferrite, where C,

N, B mix on the interstitial sublattice, while Fe, Cr, W, etc. mix on the substitutional one.

Considering a regular solution for a two-sublattice phase with A and B on the first

sublattice and C and D on the second, the excess Gibbs energy of mixing is written:

0

,:*221

*:,21

DCDCBABAxsmix LyyLyyG += (4.12)

Methods

32

where 1*:,BAL and 0

,:* DCL are regular solution parameters mixing on the sublatice

irrespective of site occupation of the other sublattice. The mole fractions (e.g. XA) are

now replaced by the occupied site fractions ( 1Ay , where 1 denotes the first sublattice). A

sub-regular model is introduced by making the interactions dependent on the site

occupation of the other sublattice as:

0

,:1220

,:1220

:,2110

:,211

DCBBACDCAAACDBADBCCBACBAxsmix LyyyLyyyLyyyLyyyG +++= (4.13)

The temperature dependence is obtained writing the parameters 0:, CBAL as polynomials of

T and lnT. The coefficients of these polynomials are stored in the SGTE database.

Fig. 4-2: Simple body-centred cubic structure with preferential occupation of atoms in the body-centre and

corner positions.

All thermodynamic calculations were carried out with the ThermoCalc database TCFe6

[123]. This database contains a new definition of the Z-phase. The limitations of the

TCFe6 database are a lack on the definition of two important alloy elements such as B

and Ta. In previous investigations by [40] a non-public database containing Ta was used,

showing qualitative different phase diagrams. Although both Ta and Nb are MX forming

elements, the atomic weight of Ta is almost twice that of Nb. Nevertheless, for a first

approximation of the phase field equilibria, thermodynamic calculations were carried out

with Nb replacing Ta. This assumption was validated by experimental results for the

alloys investigated in this work.

Methods

33

4.2. Material preparation

In the present study six 9-12% Cr steels with varying chemical compositions were

produced by vacuum induction melting with masses of about 1 kg and 4 kg at MPIE.

The production of the melts was monitored by an OBLF spark spectrometer in order to

adjust the alloy compositions.

The 12% Cr steels (masses ~ 1 kg) were hot forged in a swaging machine for rods of 4-32

mm of diameter.

The 9% Cr steels (masses ~ 4 kg) were hot rolled in a 120 mm roller mill (roll power up

to max 10 ton).

The austenisation and tempering heat treatments for all samples were carried out in

electric furnaces within a protective atmosphere of argon.

4.3. Optical microscopy characterisation

In order to study the microstructure of the different creep resistant steels, samples were

analysed in the initial state (after tempering) by light optical microscopy (LM).

Specimens were prepared by mechanical grinding (down to 1,200 grade paper), followed

by mechanical polishing with 6, 3 and 1µm diamond paste. The samples were etched

using V2A (47.5 vol% distilled water, 47.5 vol% hydrochloric acid, 4.8 vol% nitric acid

and 0.2 vol% Vogel’s reagent) for up to 50 seconds at 50°C. Optical micrographs were

taken using an Olympus BX60M microscope equipped with a Nikon DXM1200 digital

camera.

The prior austenite grain sizes (PAGS) were measured on light microscopy images using

the line intersection method. Several micrographs were analysed for each alloy to ensure

that measurements were representative for the whole material. An array of about 8

Methods

34

horizontal lines was superposed to each micrograph in order to measure the grain width

using the AnalySIS 5.0/Olympus soft imaging editor software. The results were reported

as average values from all measurements together with the error of the average value.

4.4. Transmission electron microscopy characterisation

The transmission electron microscope Tecnai Supertwin F20 with field emission gun

operating at 200kV was used in order to investigate the microstructure features which

cannot be identified using LM or scanning electron microscopy, such as sub-grains,

dislocation and nano-sized precipitates, at the initial state and after creep.

Discs of about 1.0 mm thickness were cut from creep samples and then thinned down to

0.09 mm by silicon carbide paper (1200 grade). The thinning of the specimen is very

important to minimise magnetic aberrations in the TEM [124]. The discs were twin jet

electropolished (TenuPol-5 of Struers) with a solution of acetic acid and perchloric acid

as electrolyte (95 vol% acetic acid, 5 vol% perchloric acid) at 15°C and 43 V. After

electro-polishing, the thin foils were carefully cleaned in methanol and then dried. Bright-

field micrographs and diffraction patterns were taken with a CCD camera (Gatan US

1000).

• Quantitative determination of precipitates

Observations in TEM were carried out in bright field and scanning mode (STEM) in

order to study the precipitates. Several micrographs from the sample section were

analysed for each alloy to ensure that measurements were representative for the whole

material. Precipitates were identified by a combination of electron diffraction patterns

(DF) and energy dispersive spectroscopy (EDS) analysis, to avoid ambiguous

identification of similar precipitates.

Equivalent circle diameters of the particles were calculated with an image analyser

software in order to carry out a quantitative analysis of precipitates [125]. More than 120

Methods

35

particles for each type of precipitates were quantified for reliability of the measurements.

In the case that the precipitates present a non spherical form, two perpendicular axes were

measured (a and b) and an average diameter d = (a+b)/2 was calculated.

For all the samples investigated the error was determined by Skderror ⋅±= 1 ; where S

is the standard deviation, n

k 96.11 = and n is the number of measured precipitates.

• Quantitative determination of dislocation density

The quantitative measurement of dislocation density was carried out by scanning electron

microscopy. STEM in combination with high-angle annular dark-field (HAADF) detector

is a reliable method for measurement of dislocation density in complex engineering

materials such as 9-12% Cr steels according to Pešička et al [126]. The STEM-HAADF

has traditionally been associated with good chemical contrast, related to the higher

inelastic scattering potential of heavier elements. The heavier chemical elements scatter

electrons to higher angles (~50-200 mrad), which are collected by the HAADF detector

[127] (see Fig. 4-3).

Fig. 4-3: Scheme of bright field, annular dark field and high-angle annular dark field detectors of a STEM.

Methods

36

Nevertheless, the bright contrast presented by the dislocations is associated with both de-

channeling contrast and diffuse scattering. In Refs. [128,129] it was explained that the

de-channeling contrast and the diffuse scatter are produced from static displacement of

the atoms around the dislocation core.

Additionally, favorable gr vectors can be adjusted to a two-beam case in bright field

mode (BF) by tilting the sample [127], taking into consideration that the dislocations are

invisible if the gr vector adjusted is perpendicular to the burgers vector ( br

) of the

dislocations (the so called 0=⋅bgrr or invisibility criterion). Moreover, in the multibeam

case, by aligning the main beam with a low index zone axis all the dislocations are visible

for those gr vectors which are intersected in the zone-axis that are not perpendicular to

the burgers vector (br

) of the dislocations. By adjusting a multibeam case in a zone of

low index in combination with STEM-HAADF, it is possible to have a very good

dislocation contrast, which allows the observation of dislocation structures with a higher

quality than traditional bright field or dark field TEM [126,128, 130].

In order to ensure reliable dislocation density measurements, at least 6 different

micrographs, containing between two to four sub-grains were investigated for each

sample. For each sub-grain a multiple beam case in a zone of low index using Bragg lines

was adjusted, thus highlighting most of the dislocations.

Methods

37

In Fig. 4-4 three different multibeam cases for

a particular sub-grain are shown,

corresponding to [131], [110] and [531]. As

observed the contrast for identification of

dislocations is best for the low index case [110]

(compare index adjustments [531] and [110]).

Once the area for quantification was adjusted, a

grid consisting of horizontal and vertical lines

was superimposed in the center of the sub-

grains. The numbers ny and nx of intersections

of dislocations with the vertical (Ly) and

horizontal (Lx) grid lines were counted [131].

The foil thickness (t) varied between 170 and

230 nm.

Fig. 4-4: Contrast of dislocations for multi-beam case in different zone axes (sample 12Cr4CoWTa-780 NHD

at initial stage). For the zone axes [131] (A) and [531] (A) a reduced contrast is obtained. For the multi-beam

case of low index [110] (B) high contrast is achieved and dislocations are highlighted.

Methods

38

The dislocation density was then calculated by the following expression [132]:

⎟⎟⎠

⎞⎜⎜⎝

⎛+=∑∑

∑∑

x

x

y

y

Ln

Ln

t1ρ (4.14)

• Quantitative determination of sub-grain size

The determination of the sub-grain size at initial state and after creep was carried out

using the line intersection method. Several STEM micrographs were taken through the

sample to obtain representative measurements [109,133]. An array of 6 reference lines,

perpendicular to the direction of the elongated sub-grains was set for each micrograph to

measure the sub-grain widths using the AnalySIS 5.0/Olympus soft imaging editor

software. The results were reported as average values from all measurements together

with the error of the average value.

4.5. Mechanical testing

• Creep tests

Tensile creep tests in air at constant temperature of 650°C (± 5 K) with constant load

between 80 and 250 MPa were carried out to determine the creep rupture times. Standard

cylindrical samples according to DIN50125 B 4x20 were used with 40 mm gauge length

and 4 mm diameter.

• Hardness

Vickers hardness measurements (HV10) according to DIN EN ISO 6507-1 were carried

out using a Universal Wolpert macrohardness DIA Testor 2n. At least 10 measurements

were performed on the samples in the initial state (after tempering) and after creep (to

rupture).

Results of 12% Cr Steels

39

5. Results of 12% Cr steels

For better understanding of the results, this chapter was separated in two sections:

The first section 5.1. deals with the design of two sets of martensitic/ferritic 12% Cr

alloys supported by thermodynamic modeling. The alloys were produced and creep tested

at 650°C / 80-250 MPa to investigate the microstructure evolution and the mechanical

properties.

The first group (alloy 12Cr4CoWTa) was designed in order to have a microstructure with

MX and Laves phase in the temperature range of interest (650°C), whereas the second

group (alloy 12Cr2CoWV) was designed to have a microstructure with Laves phase, MX

and M23C6 precipitates in the investigated range of temperatures.

A detailed characterisation of the microstructure of the alloys at initial state (after

tempering) and the different creep times (to rupture) was carried out by STEM.

The results of the microstructure analysis were related to the observed creep behaviour in

order to investigate the influence of the MX precipitates combined with Laves phase

(12Cr4CoWTa), and MX precipitates combined with M23C6 carbides and Laves phase

(12Cr2CoWV) on the creep strength.

The second section 5.2. focuses on the influence of processing parameters (hot

deformation and tempering temperature) on the microstructure evolution and creep

strength of one of the designed alloys (12Cr4CoWTa). The investigations were focused

on the quantitative determination of dislocation density and the sub-grain size evolution

by STEM-HAADF during creep at 650°C in the early stages of creep (4,000 h) and their

correlation with the creep test results.


40

5.1 Alloy design and characterisation

5.1.1. Alloy design of 12% Cr steels supported by ThermoCalc

The design of the novel 12% Cr alloys is based on previous work of Sauthoff et al. [40].

In a first step, ThermoCalc calculations were carried out for a reference alloy of

composition 12% Cr, 0.2% Si, 0.2% V, 0.1% C and 0.05% N to determine the influence

of the main alloy elements on the austenite stability and on the Laves phase formation.

• Influence of Co and W on microstructure formation

Influence of Co content

High contents of Co were intentionally used in order to stabilise the austenitic field, as

reported in Ref. [40]. Co shows a high solubility in the ferrite and low solubility in the

precipitates, and hence Co remains in the matrix as a solid solution. As an example, the

influence of Co on the austenite stability of the reference alloy calculated by ThermoCalc

is shown in Fig. 5-1. A single-phase austenite field above 1055°C is found for high Co

contents above 1% (in order to avoid the formation of δ-ferrite in the reference alloy).

Influence of W content

W is well known to increase the high temperature strength via solid solution hardening,

suppressing the recovery of the martensitic matrix and increasing the stability of the

precipitates by decreasing the self-diffusion rate [71,72]. W is the most potent Laves

phase former [40] and for this reason it is interesting to determine the necessary amount

of W in the reference alloy in order to form Laves phase at the temperature of interest

(650°C). In Fig. 5-2 it can be seen that the Laves phase precipitates for more than 1% W

at 650°C. The amount of W is restricted to 5.5% in order to obtain single-phase austenite

field at temperatures above 1085°C, avoiding the formation of δ-ferrite which decreases

the creep strength. In the phase field of interest, also ferrite and M23C6 carbides are


41

observed. The M23C6 precipitate is of the type (Cr,W,Fe)23C6 whereas the Laves phase is

of the type Fe2W.

The calculations also predicted a stable Z-Phase of the type CrVN, what is indeed

expected due to the high amount of Cr in the alloy [92]. A high Cr content increases the

driving force for the precipitation of the Z-Phase, which is more stable compared to V-

MX carbonitride in 12% Cr alloys [50].


42

Fig. 5-1: ThermoCalc phase fields as a function of the Co content for the reference alloy (F=ferrite and A=

austenite, ThermoCalc TCFe6).

Fig. 5-2: ThermoCalc phase fields as a function of the W content for the reference alloy (F=ferrite and A=

austenite, ThermoCalc TCFe6).


43

• 12Cr4CoWTa and 12Cr2CoWV design

Alloy 12Cr4CoWTa

The chemical composition of alloy 12Cr4CoWTa was designed in order to produce a

microstructure with ferrite, MX carbonitrides, Laves phase and Cu precipitates without

M23C6 carbides at the temperature investigated (650°C).

Based on the ThermoCalc modeling and previous observations of Schneider and Inden

[121], alloy 12Cr4CoWTa was designed with W contents of 3.5% in order to increase the

number of nuclei for Laves phase precipitation and to decrease the critical nucleus size,

so that a finer distribution of the Laves phase is obtained.

Based on the ThermoCalc calculations, it was determined that for carbon contents of up

to 0.06% and nitrogen contents of 0.05%, the formation of carbonitrides is promoted and

the formation of M23C6 carbides is negligible. In general, the modeling showed, as

expected, that low amounts of C decrease the quantity of M23C6 precipitates, while the

volume fraction of MX carbonitrides significantly increases. Moreover, the amount of

carbon plays a significant role in the stabilisation of the austenite at high temperatures,

influencing the martensitic transformation.

The content of Ta was 0.8%. The value of Ta was taken based on previous observations

of Sauthoff et al. [40] for precipitation of Laves phase and suppression of M23C6 carbides.

The relatively high Ta content was set to promote the formation of a high volume fraction

of Ta-containing MX precipitates, despite of the well known disadvantages of high Ta

concentrations, such as the promotion of δ-ferrite formation and the precipitation of MX

particles from the liquid state, which results in difficulties to control the size and the

distribution of the precipitates, decreasing the creep strength. The B content was set to

0.001%. The addition of B stabilises the microstructure, improving the stability of the

lath martensite and hence the creep rupture strength [39]. Due to database limitations, the

influence of B on the phase formation can not be calculated. Abe investigated the BN

precipitation in 9% Cr heat resistant steels and according to his observations, BN should

not form for the combination of 0.06% N and 0.001% B. [58].


44

The ThermoCalc calculations show that alloy 12Cr4CoWTa contains a ferrite, austenite,

MX carbonitrides, Laves phase, W-Cu rich precipitates and Z-phase at the tempering

temperature (780°C) and ferrite, MX, Laves phase, Z-phase and Cu rich precipitates

without M23C6 carbides at the temperature investigated (650°C). Tab. 5-1 shows the

calculated chemical composition of the precipitates at 780°C and at 650°C. Tab. 5-3

shows the volume fraction of the phases at the temperatures of interest (780°C and 650°C,

ThermoCalc TCFe6)

Alloy 12Cr2CoWV

Alloy 12Cr2CoWV was designed in order to obtain ferrite, MX carbonitrides, M23C6

carbide, Laves phase and Cu precipitates at the temperature of interest (650°C). The W

content was set to 3.6% to promote the Laves phase formation in order to increase the

number of nuclei for Laves phase precipitation and to decrease the critical nucleus size,

so that a finer distribution of the Laves phase is obtained. 0.25% V and 0.15% Ta was

used to obtain fine dispersion of V and Ta carbonitrides. The MX precipitates proved to

be beneficial to the creep rupture strength [64] because they are extremely stable and they

show slow coarsening rate during creep. The C content was set to 0.15% and the N

content to 0.06% to promote the formation of M23C6 precipitates and MX carbonitrides

[41]. Relatively high B content (0.03%) was added in order to stabilise the M23C6

carbides [26].

The ThermoCalc calculations show that alloy 12Cr2CoWV contains ferrite, MX

carbonitrides, M23C6, Laves phase, Z-phase and W-Cu rich precipitates at the tempering

temperature (780°C) and ferrite, MX, M23C6, Laves phase, Z-phase and Cu rich

precipitates at 650°C. Tab. 5-2 shows the chemical composition of the precipitates at

780°C and at 650°C. Tab. 5-4 shows the volume fraction of the phases at the

temperatures of interest (780°C and 650°C).

The final compositions of the two alloys are given in Tab. 5-5.


45

Tab. 5-1: Calculated composition (wt%) of precipitates for alloy 12Cr4CoWTa.

T°C Phases Fe Cr W Nb C N Co Mn Cu Ferrite 78.88 12.70 2.28 - - - 4.60 0.66 0.88

Austenite 78.52 11.60 1.65 - - - 5.14 1.31 1.78 MX - 0.77 - 87.49 6.57 5.17 - - -

Laves phase 30.12 7.70 60.51 1.34 - - 0.28 0.05 - Z-Phase 7.12 26.00 - 58.32 - 8.56 - - -

780

W-Cu 0.60 - 57.91 - - - 0.12 0.38 40.99 Ferrite 80.80 12.67 0.69 - - - 4.80 0.74 0.30

MX - 0.91 - 87.61 7.97 3.51 - - - Laves phase 28.52 8.98 61.65 0.63 - - 0.13 0.09 -

Z-Phase 7.06 26.04 - 58.28 - 8.62 - - - 650

Cu 0.29 - 5.73 - - - 1.22 1.20 91.56 Tab. 5-2: Calculated composition (wt%) of precipitates for alloy 12Cr2CoWV.

T°C Phases Fe Cr W V C N Nb Co Cu

Ferrite 82.64 11.27 2.55 0.07 - - - 2.61 0.86 Laves phase 30.22 7.29 62.22 - - - 0.13 0.14 -

Z-Phase 5.36 34.13 - 23.67 - 10.16 26.68 - - M23C6 19.95 50.12 23.96 1.26 4.63 - - 0.08 -

780

W-Cu 0.59 - 58.52 - - - - 0.07 40.82 Ferrite 85.02 11.13 0.79 0.06 - - - 2.71 0.29

Laves phase 29.16 8.32 62.27 - - - 0.18 0.07 - Z-Phase 5.24 34.86 - 25.35 - 9.77 24.78 - - M23C6 14.67 60.01 18.95 1.43 4.89 - - 0.05 -

650

Cu 0.25 - 5.67 - - - - 0.59 93.49


46

Tab. 5-3: Volume fraction (%) of precipitates for alloy 12Cr4CoWTa calculated with ThermoCalc at tempering

temperature (780°C) and at the creep testing temperature (650°C) (MX are of the type Nb(C,N)).

T°C Phases Volume Fraction T°C Phases Volume

Fraction Ferrite 88.94 Ferrite 95.11

Austenite 8.80 MX 0.75 MX 0.89 Laves phase 2.78

Laves phase 1.18 Z-Phase 0.28 Z-Phase 0.04 Cu 1.08

780

W-Cu 0.15

650

- - Tab. 5-4: Volume fraction (%) of precipitates for alloy 12Cr2CoWV calculated with ThermoCalc at tempering

temperature (780°C) and at the creep testing temperature (650°C) (MX are of the type Nb(C,N)).

T°C Phases Volume Fraction T°C Phases Volume

Fraction Ferrite 95.88 Ferrite 93.28

Laves phase 0.16 Laves phase 2.27 Z-Phase 0.71 Z-Phase 0.74 M23C6 3.00 M23C6 3.03

780

W-Cu 0.25

650

Cu 0.68


47

5.1.2. Alloy production

The alloys were prepared by vacuum induction melting with masses of about 1 kg. The

final chemical composition of the alloys measured by chemical analysis is shown in Tab.

5-5.

Tab. 5-5: Analysed chemical composition (wt%) of the model alloys investigated.

Alloy 12Cr4CoWTa

Cr Mn Ta W V Cu C B N Si Co 12.9 ±0.5

0.6 ±0.1

0.85 ±0.05

3.8 ±0.2 - 1.0

±0.2 0.06

±0.01 0.001

±0.0010.05

±0.01 0.5

±0.05 4.2

±0.10

Alloy 12Cr2CoWV

Cr Mn Ta W V Cu C B N Si Co 12.6 ±0.5

0.7 ±0.1

0.16 ±0.05

3.6 ±0.2

0.24 ±0.05

1.0 ±0.2

0.15 ±0.01

0.033 ±0.005

0.06 ±0.01

0.5 ±0.05

2.5 ±0.10

Samples were hot forged at 1150°C (resulting in 70% final area reduction) with posterior

air cooling. Heat treatments were carried out based on the standard heat treatments for

12% Cr heat resistant steels [31] for power plant parts:

• Austenitisation at 1070°C for 0.5 h followed by air-cooling (with martensite

transformation).

• Tempering for 2 h at 780°C with subsequent air-cooling (with recovery of ductility,

annihilation of dislocations and precipitation of carbonitrides, M23C6 carbides or

Laves phase).

The treatments for alloys 12Cr4CoWTa and 12Cr2CoWV are shown in Fig. 5-3.


48

Fig. 5-3: 12% Cr steels heat treatment scheme.

5.1.3. Microstructure evolution (precipitates quantification)

Light optical microscopy investigations of both alloys after heat treatment show a

martensitic/ferritic matrix and precipitates. TEM and STEM investigations are carried out

to quantify the microstructure features of the alloys in the initial stage and after different

creep times.

• Alloy 12Cr4CoWTa-780 HD

Initial microstructure: In Fig. 5-4 a STEM micrograph of alloy 12Cr4CoWTa-780 HD

in the initial condition after tempering at 780°C for 2 h is shown. Two precipitates are

identified by EDS analysis: Laves phase and MX carbonitrides. A detailed TEM

investigation of the MX particles (see Fig. 5-5) showed two types of MX, relatively large

C rich Ta-MX precipitates with an average size of 137 ± 15 nm, which were not

dissolved in the austenisation treatment and N rich Ta-MX particles with an average size

of 30 ± 2 nm. The average size of Laves phase was 196 ± 35 nm in the initial stage (see

Tab. 5-6).


49

Fig. 5-4: TEM image of alloy 12Cr4CoWTa-780 HD

in the initial stage. Black arrows show MX particles,

white arrows show Laves Phase.

Fig. 5-5: TEM image of alloy 12Cr4CoWTa-780 HD

in the initial stage. C rich Ta-MX (white arrows) and

N rich Ta-MX particles (black arrows).

Tab. 5-6: Mean size of precipitates in alloy 12Cr4CoWTa-780 HD (time in hours and size in nanometers).

12Cr4CoWTa-780 HD 0 h 1,200 h 3,650 h

N rich Ta-MX 30 ± 2 31 ± 2 32 ± 2

C rich Ta-MX 137 ± 15 - 142 ± 8

Laves phase 196 ± 35 268 ± 34 302 ± 49

Microstructure after creep: After 3,650 h creep at 650°C, alloy 12Cr4CoWTa-780 HD

presented the same precipitates as in the initial stage (C rich Ta-MX, N rich Ta-MX and

Laves phase, Fig. 5-6). The average size of the Laves phase was 302 ± 49 nm. Very few

Laves phase particles with a very large size of up to 787 ± 49 nm were observed. In Fig.

5-7 and Fig. 5-8 a diffraction pattern as well as an EDS spectrum of the Laves phase after

3,650 h is shown. The Laves phase presented a hexagonal structure, the main chemical

elements were W, Fe and some of Cr. Slow coarsening of C rich Ta-MX was observed,

with an average final size of 142 ± 8 nm (137 ± 15 nm in the initial state, see Tab. 5-6).

The size of the N rich Ta-MX remains almost constant with an average size of 30 ± 2 nm.

Z-phase was not observed, probably due to short creep times.


50

Fig. 5-6: TEM image of sample 12Cr4CoWTa-780 HD after 3,650 h under creep condition 650°C at 100 MPa.

Laves phase (black arrows) and MX precipitates (white arrows) are present in the microstructure.

Fig. 5-7: Diffraction pattern of Laves phase in

sample 12Cr4CoWT-780 HD after 3,650 h creep.

Fig. 5-8: EDS spectrum of Laves phase in sample

12Cr4CoWTa-780 HD after 3,650 h creep.


51

• Alloy 12Cr2CoWV-780 HD

Initial microstructure: TEM investigations of alloy 12Cr2CoWV-780 HD after

tempering revealed M23C6 carbides, MX precipitates and nano-sized Cu precipitates, as

well as few W-Cu inclusions in the initial microstructure (Fig. 5-9). EDS analysis showed

that the main elements of the MX precipitates were V, Ta, C and N (Fig. 5-10). These

MX carbonitrides showed an average size of 20 ± 2 nm in the initial stage (Tab. 5-7). The

fcc M23C6 precipitates were of the type (Cr,Fe,W,V)23C6 (Fig. 5-11 and 5-12) and form

principally at the grain boundaries with an average size of 140 ± 10 nm.

Tab. 5-7: Mean size of precipitates in alloy 12Cr2CoWV-780 HD (time in hours and size in nanometers).

12Cr2CoWV-780 HD 0 h 2,000 h 6,150 h

(V,Ta)(C,N) 20 ± 2 36 ± 2 38 ± 2

M23C6 140 ± 10 180 ± 8 185 ± 9

Laves phase - 285 ± 42 307 ± 41

Z-Phase - - 273 ± 47

Fig. 5-9: TEM image of alloy 12Cr2CoWV-780 HD

in the initial stage. M23C6 carbides (white arrows)

and MX precipitates (black arrows).

Fig. 5-10: EDS spectrum of MX of the type

(V,Ta)(C,N) showing V and Ta as main elements.


52

Fig. 5-11: Diffraction pattern of M23C6 carbide in

the initial stage of alloy 12Cr2CoWV-780 HD (fcc

crystal structure).

Fig. 5-12: EDS spectrum M23C6 precipitate in the initial

stage of alloy 12Cr2CoWV-780 HD showing the main

elements Cr, V, W and Fe.

Microstructure after creep: Figures 5-13 to 5-15 show the microstructure of alloy

12Cr2CoWV-780 HD after 6,150h creep. M23C6 precipitates of an average size of 185 ±

9 nm and MX precipitates of the type (V,Ta)(C,N) with an average size of 38 ± 2 nm

were observed. After 2,000 h under creep conditions Laves phase was observed. At 6,150

h the average particle size of Laves phase was 307 ± 41 nm. After 6,150 h the Z-phase

was observed in the microstructure of alloy 12Cr2CoWV-780 HD (Fig. 5-15 and Fig. 5-

16). EDS analysis revealed that Cr, Ta, V and N were the main elements in this phase.


53


after 6,150 h creep. M23C6 carbide (white arrows) and

Laves phase (black arrows) are observed.


after 6,150 h creep showing a nano-sized MX particle

of the type (V,Ta)(C,N) (white arrow).

Fig. 5-15: STEM image of alloy 12Cr2CoWV-780

HD after 6,150 h creep showing Laves phase and

M23C6 carbides (black arrows) and Z-Phase (white

arrow). The white points in the Z-phase indicate the

EDS measurements.

Fig. 5-16: EDS of Z-phase particle shown in Fig. 5-15.

Three measurements were carried out in this phase

(white points). The main elements are V, Cr, Ta and

N.


54

• Creep results

Results of the creep tests are shown in Fig. 5-17. Creep tests at 650°C showed an

increment of the creep strength for alloy 12Cr2CoWV-780 HD compared to alloy

12Cr4CoWTa-780 HD. Both alloys presented higher creep strength at high tensile load

(175-250 MPa) compared to the reference alloy P92, but a decrease in the creep strength

was observed for long-term creep at relatively low tensile loads (80-150 MPa). This

behaviour was more pronounced in alloys 12Cr4CoWTa-780 HD.

Fig. 5-17: Results of the tensile creep tests showing times to rupture as a function of applied stress for alloys

12Cr4CoWTa-780 HD and 12Cr2CoWV-780 HD at 650°C. An increased creep strength for alloy

12Cr2CoWV-780 HD can be observed. Results of creep tests of a P92 steel under similar conditions [42] are

shown as reference.


55

5.2. Influence of processing parameters

In the present section, alloy 12Cr4CoWTa was selected in order to investigate the

influence of the hot deformation and tempering temperature on the microstructure

evolution and creep strength. The investigations were focused on the quantitative

determination of dislocation density and the sub-grain size evolution by STEM-HAADF

during creep at 650°C.

5.2.1. Alloy processing

In order to investigate the influence of the hot deformation process two samples were

prepared as follows:

The first sample (12Cr4CoWTa-780 HD) was hot forged, austenitised and tempered with

following parameters:

• Hot forging at 1150°C with subsequent air cooling (70% final area reduction / hot

forging for homogenisation of the material).

• Austenisation at 1070°C for 0.5 h followed by air-cooling (with martensite

transformation).

• Tempering at 780°C for 2 h with subsequent air-cooling (with recovery of ductility,

annihilation of dislocations and precipitation of carbonitrides and Laves phase).

The second sample (12Cr4CoWTa-780 NHD) was austenitised and tempered as sample

12Cr4CoWTa-780 HD, but the hot-deformation process was not carried out (NHD).

In order to study the influence of the tempering temperature, a third sample

(12Cr4CoWTa-680 HD) was prepared as follows:

• Hot forging at 1150°C with subsequent air cooling.

• Austenisation at 1070°C for 0.5 h followed by air-cooling.


56

• Tempering at 680°C for 2 h with subsequent air-cooling.

The tempering temperature of 680°C was chosen because it is a typical heat treatment

temperature for components that require high tensile strength.

The various treatments for alloys 12Cr4CoWTa are shown in Fig. 5-18.

Fig. 5-18: Alloy 12Cr4CoWTa heat treatment scheme.

5.2.2. Initial microstructure after tempering (dislocation density and sub-grain size)

Samples prepared from alloy 12Cr4CoWTa with different processing parameters

presented a martensitic/ferritic matrix with a relatively high dislocation density after

tempering. Several types of internal interfaces as prior austenite grain boundaries, prior

packet or block boundaries, prior martensite lath boundaries and sub-grain boundaries are

observed in all samples (see Fig. 5-19). Nucleation of most carbides is heterogeneous and

preferentially takes place at prior austenite grain boundaries and prior packet or block

boundaries.

As an example, a HAADF-STEM image of the initial microstructure of alloy

12Cr4CoWTa-780 HD is shown in Fig. 5-20. Free dislocations and sub-grains inside of a

prior martensite lath are seen. Laves phase precipitated at prior austenite grain boundaries

and lath or block boundaries and MX carbonitrides precipitated inside the sub-grains or at

the sub-grain boundaries can also be clearly observed.


57

Fig. 5-19: STEM-HAADF image of the initial microstructure of sample 12Cr4CoWTa-780 HD (inverse

contrast). The square area was amplified for better observation of the internal interfaces. A prior austenite grain

boundary (dashed line), prior martensite laths (dotted lines) and sub-grain boundaries (full lines) are shown.

Fig. 5-20: Montage of STEM-HAADF images of the initial microstructure of sample 12Cr4CoWTa-780 HD

(inverse contrast). Precipitates, sub-grains and dislocations are observed. For each picture of the montage, a

multi-beam with a low index zone axis was adjusted in order to highlight the dislocation inside the sub-grains.


58

5.2.3. Microstructure features at the initial state

The measurements of prior austenite grain size (PAGS), sub-grain size, dislocation

density and hardness at initial state (after tempering) are shown in Tab. 5-8. Hot

deformed samples showed similar PAGS values. For example, in 12Cr4CoWTa-780 HD

the PAGS was 33 ± 3 μm and 32 ± 3 μm for 12Cr4CoWTa-680 HD. The 12Cr4CoWTa-

780 NHD showed the largest PAGS value (139 ± 9 μm).

The sub-grain sizes measurements were very similar for all samples. 12Cr4CoWTa-780

NHD showed a sub-grain size of 450 ± 47 nm, for sample 12Cr4CoWTa-780 HD the

sub-grain size was 420 ± 32 nm, whereas samples 12Cr4CoWTa-680 HD showed a sub-

grain size of 380 ± 30 nm (see Tab. 5-8).

All samples presented a relatively high dislocation density after tempering. The sample

12Cr4CoWTa-680 HD showed the highest dislocation density with 27.3 x 1013 m-2,

followed by sample 12Cr4CoWTa-780 HD with 26.2 x 1013 m-2, whereas sample

12Cr4CoWTa-780 NHD showed the smallest value (21.2 x 1013 m-2).

The hardness is closely related to the dislocation density and sub-grain size. The sample

12Cr4CoWTa-680 HD showed the highest hardness 393 ± 1 HV10, for sample

12Cr4CoWTa-780 HD the hardness was 388 ± 1 HV10, whereas sample 12Cr4CoWTa-

780 NHD showed the smallest value (376 ± 3 HV10).

Tab. 5-8: Quantitative determination of PAGS, sub-grain size, dislocation density and hardness at the initial stage.

Samples PAGS (μm) Sub-grain size (nm)

Dislocation density ρ (1013 m-2)

Scatter in ρ (1013 m-2) HV10

12Cr4CoWTa-780 NHD initial 139 ± 9 450 ± 47 21.2 26.2 - 13.7 376 ± 3

12Cr4CoWTa-780 HD initial 33 ± 3 420 ± 32 26.2 34.8 - 13.8 388 ± 1

12Cr4CoWTa-680 HD initial 32 ± 3 380 ± 30 27.3 36.2 - 21.1 393 ± 1


59

5.2.4. Microstructure evolution during creep (dislocation density and sub-grain size)

During creep, dislocations interact with each other, with internal interfaces, with solute

atoms and with precipitated particles [103]. In Fig. 5-21 an HAADF-STEM image is

shown where dislocations interact with each other and with precipitates in sample

12Cr4CoWTa-680 HD after 2,875 h creep.

Fig. 5-21: STEM-HAADF image of sample 12Cr4CoWTa-680 HD after creep (2,875 h / 80MPa) showing the

interaction of dislocations as well as with precipitates (A). Same image with inverse contrast for better

observation of dislocation networks (B).

The working temperature (650°C) and the exposure at stress during service promote

changes in dislocation density, sub-grain size and precipitate size reducing the creep

strength [109].

The evolution of the dislocation density and the sub-grain size can clearly be seen in

sample 12Cr4CoWTa-780 HD. The dislocation density was 26.2 x 1013 m-2 in the initial

state (Tab. 5-8). After 1,121 h under creep conditions (650°C) the dislocation density

decreased to 4.7 x 1013 m-2 (Tab. 5-9) and further decreased to 2.9 x 1013 m-2 after 3,650 h

creep at 650°C (Tab. 5-10). Due to the extensive creep deformation, the original sub-

grain structure is replaced by polygonal sub-grains, where the sub-grain size changes

from 420 ± 32 to 700 ± 88 nm after 1,121 h / 650°C creep (Fig. 5-22a) and to 900 ± 86

nm after 3,650 h at 650°C (Fig. 5-22b).


60

Fig. 5-22: Montage of STEM-HAADF images for (A) sample 12Cr4CoWTa-780 HD after creep (1,121 h /

145MPa / 650°C) and (B) sample 12Cr4CoWTa-780 HD after creep (3,650 h / 80MPa / 650°C). Smaller sub-

grain sizes and higher dislocation densities are observed for the sample with shorter creep time (A). White

arrows indicate the size of some sub-grains for comparison.


61

In Fig. 5-23 and Fig. 5-24 the tensile creep curves of samples 12Cr4CoWTa-780 HD,

12Cr4CoWTa-780 NHD and 12Cr4CoWTa-680 HD are shown. Independently of the

processing parameters all curves present similar characteristics. In the creep rate vs. strain

curves, there is a short primary stage with subsequent secondary stage corresponding to a

creep rate minimum. The creep rate minimum is immediately followed by a tertiary stage

with an increment of the creep rate.

5.2.5. Influence of hot deformation on creep strength

The influence of hot deformation on creep strength for alloy 12Cr4CoWTa-780 is shown

in Fig. 5-23. As shown in Fig. 5-23a, sample 12Cr4CoWTa-780 HD presented a rupture

time of 1,121 h, more than twice the rupture time of the non hot-deformed sample

(12Cr4CoWTa-780 NHD, 503 h). Sample 12Cr4CoWTa-780 HD showed a reduced final

strain (17.6 %) compared to sample 12Cr4CoWTa-780 NHD (37.9 %).

As for the creep curves, the dislocation density measured after creep for both samples

also showed remarkable differences. In sample 12Cr4CoWTa-780 HD, a dislocation

density of 4.7 x 1013 m-2 was measured, which was almost twice the value obtained for

the non hot-deformed sample (12Cr4CoWTa-780 NHD = 2.5 x 1013 m-2, see Tab. 5-9).

Moreover, the sub-grain size of the hot-deformed sample 12Cr4CoWTa-780 HD was 700

± 88 nm, whereas a sub-grain size of 870 ± 122 nm was measured for sample

12Cr4CoWTa-780 NHD. The dislocation density and sub-grain size were directly

correlated with the hardness, being larger for the hot-deformed sample (12Cr4CoWTa-

780 HD = HV10 298 ± 1; 12Cr4CoWTa-780 NHD = HV10 262 ± 2).

In Fig. 5-23b the minimum creep rates are shown. Sample 12Cr4CoWTa-780 HD

presented a lower minimum creep rate of 1.69 x 10-8 h-1 at smaller strain (1.61%);

whereas sample 12Cr4CoWTa-780 NHD had a minimum creep rate of 4.97 x 10-8 h-1 at

6.3% strain.


62

Fig. 5-23: Tensile creep test curves comparing the creep strength of sample 12Cr4CoWTa-780 NHD and

sample 12Cr4CoWTa-780HD at 145 MPa / 650°C. (A) strain vs. time (B) creep rate vs. strain.

Tab. 5-9 Effect of hot-deformation: Comparison of sub-grain size, dislocation density and hardness for samples

12Cr4CoWTa-780 NHD and 12Cr4CoWTa-780 HD after creep.

Samples Stress (MPa)

Rupture time (h)

Sub-grain size (nm)



12Cr4CoWTa-780 NHD 145 503 870 ± 122 2.5 5.4 - 1.2 262 ± 2

12Cr4CoWTa-780 HD 145 1,121 700 ± 88 4.7 6.3 - 1.8 298 ± 1

5.2.6. Influence of tempering temperature on creep strength

Samples 12Cr4CoWTa-680 HD and 12Cr4CoWTa-780 HD were tested under the same

creep conditions. For both samples, the microstructure after creep showed large sub-

grains with low dislocation density in the interior. For sample 12Cr4CoWTa-680 HD the

dislocation density decreased from 27.3 x 1013 m-2 in the initial state to 2.7 x 1013 m-2

after creep (compare Tab. 5-8 and Tab. 5-10). For the sample tempered at 780°C, the

dislocation density decreased from 26.2 x 1013 m-2 to 2.9 x 1013 m-2. In addition the

minimum creep rate for both samples was very similar (1.24 x 10-9 h-1 at 0.88% strain for

12Cr4CoWTa-680 HD and 8.56 x 10-10 h-1 at 0.76% strain for sample 12Cr4CoWTa-780

HD). The creep curves in Fig. 5-24a show very similar creep behaviour for both samples.

The difference lies in the higher slope presented in sample 12Cr4CoWTa-680 HD, which


63

leads to a shorter rupture time of 2,875 h compared to sample 12Cr4CoWTa-780 HD

with a rupture time of 3,650 h. Sample 12Cr4CoWTa-680 HD presented a reduced

ductility (strain 11.9%) compared to sample 12Cr4CoWTa-780 HD (strain 17.4%). Fig.

5-24b shows the differences in strain at rupture. Remarkable differences were observed in

the sub-grain size evolution during creep. For sample 12Cr4CoWTa-680 HD the sub-

grain size increased from 380 ± 30 nm (initial stage) to 670 ± 72 nm (after creep).

For sample 12Cr4CoWTa-780 HD, the sub-grain size increased from 420 ± 32 nm to 900

± 86 nm. A comparison of the sub-grain size after creep is seen in Fig. 5-25. Sample

12Cr4CoWTa-680 HD (Fig. 5-25a) showed a smaller sub-grain size than sample

12Cr4CoWTa-780 HD (Fig. 5-25b).

The hardness measurements showed higher values for 12Cr4CoWTa-680 HD than for

12Cr4CoWTa-780 HD. For sample 12Cr4CoWTa-680 HD the hardness value was 296 ±

1 HV10, whereas for samples 12Cr4CoWTa-780 HD the hardness was 279 ± 2 HV10 (Tab.

5-10).

Fig. 5-24: Tensile creep test curves comparing sample 12Cr4CoWTa-780 HD and sample 12Cr4CoWTa-680

HD at 80 MPa / 650°C. (A) Strain vs. time (B) creep rate vs. strain.


64

Tab. 5-10 Effect of tempering temperature: Comparison of sub-grain size, dislocation density and hardness for

samples 12Cr4CoWTa-680 HD and 12Cr4CoWTa-780 HD after creep.

Samples Stress (MPa)

Rupture time (h)

Sub-grain size (nm)



12Cr4CoWTa-680 HD 80 2,875 670 ± 72 2.7 5.8 - 1.1 296 ± 1

12Cr4CoWTa-780 HD 80 3,650 900 ± 86 2.9 6.0 - 1.2 279 ± 2

Fig. 5-25: STEM-HAADF image of (A) sample 12Cr4CoWTa-680 HD after creep (2,875 h) and (B) sample

12Cr4CoWTa-780 HD after creep (3,650 h). Smaller sub-grain sizes are observed on sample (A).

In Fig. 5-26 results of several creep rupture tests are shown for both samples at different

loads at 650°C. It was observed that the sample 12Cr4CoWTa-680 HD presented higher

creep strengths than 12Cr4CoWTa-780 HD only after short-term creep with higher

stresses.


65

Fig. 5-26: Time to rupture as a function of applied tensile stress for 12Cr4CoWTa-680 HD and 12Cr4CoWTa-

780 HD.

Discussion of 12% Cr Steels

66

6. Discussion of 12% Cr steels

6.1 Alloy design and characterisation

6.1.1. Microstructure evolution

• Alloy 12Cr4CoWTa-780 HD

Initial microstructure: It is important to note that for sample 12Cr4CoWTa-780 HD, the

precipitation of Laves phase occurs during tempering at 780°C (see Fig. 5-4).

The microstructure observations (Fig. 5-5) suggest that the C rich Ta-MX forms in the

liquid state and they are present as undissolved particles during the austenisation.

Previous investigations have shown the formation of similar primary C rich Nb-MX in

the as normalised condition on 9Cr1MoVNb steel [134]. Nevertheless, both C rich Ta-

MX and the Laves phase, are usually found on prior austenite grain boundaries and lath

or block boundaries. At this sites the effective surface energy is lower, thus diminished

the free energy barrier and facilitating nucleation [114], whereas the N rich Ta-MX form

more homogenously within sub-grains and at sub-grains boundaries, as described in Ref.

[82]. The ThermoCalc calculations at equilibrium conditions at this temperature predict

the presence of Laves phase, a carbon rich MX, Z-phase and W-Cu precipitates (Tab. 5-1

/ 780°C). The thermodynamic calculations showed good agreement with the

microstructure observation, except for the W-Cu particles and Z-phase. Predicted W-Cu

particles were not observed, probably due to the very low volume fraction of this phase

(Tab 5-3 / 780°C). Z-phase was not observed in the initial microstructure due to the slow

kinetics of the Z-phase precipitation. In Ref. [92] it was demonstrated that N-rich MX

precipitates are gradually transformed into Z-phase after 10,000-30,000 h.

Microstructure after creep: After 3,650 h creep at 650°C, the evolution of the

precipitates in the microstructure can be observed. Especially the distribution of the

Laves phase is not homogeneous, showing small Laves phase particles of 70 ± 35 nm

throughout the microstructure. The relatively small size of the Laves phase compared to


67

the size of the Laves phase observed in 12%Cr-Mo steels under similar creep conditions

[135] may be related to the high W content in this alloy (3.8%), which produces a large

number of finely dispersed nuclei for Laves phase formation [40].

The C rich Ta-MX and the Laves phase usually precipitate on the prior austenite grain

boundaries and lath or block boundaries, whereas the N rich Ta-MX forms within the

sub-grains or at sub-grains boundaries. The C rich Ta-MX and the Laves phase may

pinned the prior austenite grain boundaries until the average particle size increases with

time due to the coarsening process.

The primary C rich Ta-MX showed a larger size (142 ± 8 nm) compared with N rich Ta-

MX (32 ± 2 nm). As was previously explained C rich Ta-MX precipitates are present as

undissolved particles during the austenisation, due to their high stability, whereas the

secondary N rich Ta-MX form during tempering at 780°C [42]. The ThermoCalc

calculations showed good agreement with the microstructure observation except for the

Z-phase. Z-phase was not observed, probably due to short creep times. Cipolla [19]

demonstrated that MX particles converted into modified Z-Phase via uptake of Cr from

the matrix followed by crystallographic transformation of the crystal structure from cubic

to tetragonal at long creep times. So we can conclude that the N rich MX observed in the

initial microstructure will be transformed into Z-Phase after several thousand hours. No

BN was observed in any sample of alloy 12Cr4CoWTa-780 HD.

• Alloy 12Cr2CoWV-780 HD

Initial microstructure: The ThermoCalc calculations predict the presence of Laves

phase, M23C6 carbides, Z-phase and W-Cu precipitates at 780°C (see Tab. 5-2 / 780°C).

The ThermoCalc calculations are in good agreement whit the microstructural

investigation, except for Laves phase and Z-phase. As expected, no Laves phase was

present in the microstructure of alloy 12Cr2CoWV-780 HD after tempering, due to the

competitive growth between M23C6 carbides and the Laves phase. The tempering

treatment led to the formation of M23C6, while the amount of Laves phase was negligible


68

at this stage. The reduction in temperature from 780 to 650°C after the tempering process

led to a clear reduction in the growth rate of M23C6. This is due to the reduced mobilities

of the rate controlling elements W and Cr [85]. The driving forces for diffusion from the

ferritic matrix to the precipitates were increased by the change in temperature [85]. The

metallurgical reasons for this slow Laves phase growth may be explained by the slow

mobility of the rate controlling element W in this phase [85].

Laves phase precipitates during creep (tensile creep test with constant load at 650°C) on

prior austenite grain boundaries and lath or block boundaries [13] as well as on existing

Cu-precipitates [53] due to the lower surface energy, thus the free energy barrier

decreases facilitating nucleation. As for alloy 12Cr4CoWTa-780 HD, Z-phase is shown

as a stable phase in alloy 12Cr2CoWV-780 HD in the equilibrium condition. As

explained for alloy 12Cr4CoWTa-780 HD, no Z-phase is expected in the initial

microstructure, due to the nucleation kinetics of the Z-phase, which precipitates later on

existing MX carbonitrides during creep.

Microstructure after creep: After 6,150 h the Z-phase was first observed in the

microstructure of alloy 12Cr2CoWV-780 HD. This is in accordance with previous

observations of Danielsen and Hald in Ta-rich heat resistant steels, where a Z-phase of

the type Cr(V,Ta)N formed [89]. The volume fraction of the Z-phase was very low (see

Tab. 5-4) and only a very small amount of Z-phase was identified. After 6,150 h creep,

the phases observed were in agreement with the equilibrium phases calculated with

ThermoCalc for equilibrium conditions at 650°C, except for the MX carbonitride, which

is still present in the microstructure after 6,150 h. We can expect that by reaching the

equilibrium conditions almost all MX carbonitrides should have transformed into Z-phase

following the mechanism explained by Cipolla in [19]. BN was not observed in any

sample of alloy 12Cr2CoWV-780 HD. The MX carbonitrides showed slow growth and

coarsening rates [136].


69

• Creep

The increased creep strength of alloy 12Cr2CoWV-780 HD may be related to the

observed and calculated formation of numerous dispersed nano-sized MX carbonitrides,

which are mainly responsible for the pinning of dislocations and sub-grains during creep.

In alloy 12Cr4CoWTa-780 HD two types of MX were observed, relatively large C rich

Ta-MX with a final mean size of 142 ± 8 nm and a volume fraction of 0.75% (according

to ThermoCalc calculations), as well as nano-sized N rich Ta-MX with a final mean size

of 32 ± 2 nm and a volume fraction of 0.28%. The presence of large C rich Ta-MX could

be a reason for a reduced pinning of dislocation in alloy 12Cr4CoWTa-780 HD compared

to 12Cr2CoWV-780 HD, where a 0.74% volume fraction of nano dispersed (V,Ta)(C,N)

with a final mean size of 38 ± 2 nm is found; but this fact is very difficult to be proved,

because a combination of several strengthening mechanisms is present during creep

[120,137].

Another reason for the high creep strength of alloy 12Cr2CoWV-780 HD may be

explained by the presence of the M23C6 precipitates. M23C6 carbides may improve the

creep strength due to the relatively high volume fraction of this phase (3.03% vol.

fraction, Tab. 5-4 / 650°C), which provides a higher mechanical stability to the alloy. A

similar observation was recently reported by Aghajani et al. [135] in 12% Cr tempered

martensitic/ferritic steels after long-term creep of 139,971 h. Moreover, the effect of B

should also be taken into consideration, especially in the presence of M23C6 carbides. As

reported by Abe [26], B decreases the Oswald ripening rate of M23C6 carbides by an

enrichment of B in the vicinity of prior austenite grains.

The enrichment process of boron in M23C6 is schematically shown in Fig. 6-1. First, in

the austenisation step, the segregation of boron takes place at grain boundaries. During

subsequent tempering, precipitation of M23C6 carbides takes place preferentially at prior

austenite grain boundaries and lath or block boundaries. Because of the segregation of

boron at grain boundaries the enrichment of boron in M23C6 to form M23(C,B)6 occurs in

the vicinity of prior austenite boundaries [39].


70

Fig. 6-1: Formation process of M23(C,B)6 during heat treatment [39].

The stabilisation of the fine M23C6 carbides may retard the tertiary creep and increase the

time to rupture [72]. B and N additions should be balanced carefully to avoid the

formation of large BN particles, which are detrimental for the creep strength [58].

6.2. Influence of processing parameters

6.2.1. Initial microstructure after tempering (dislocation density and sub-grain size)

In general the high dislocation density after tempering, as explained in point 5.5, is

produced by the martensitic transformation during air-cooling from the austenisation heat

treatment temperature [132]. The transformation causes a large local deformation of the

matrix, resulting in strong work hardening due to a cellular dislocation structure with

high dislocation density. Tempering of the material allows the recovery of the martensitic

structure with transformation into a ferritic structure. Moreover the tempering leads to

precipitation of solute atoms and to recovery of the dislocation cell structure resulting in a

sub-grain structure [37]. Nucleation of most carbides is heterogeneous and preferentially

takes place at prior austenite grain boundaries and prior packet or block boundaries.

During tempering, MX carbonitrides precipitate more homogeneously inside the sub-

grains, probably due to lattice coherence, facilitating nucleation [42]. They remain

distinctly smaller than the heterogeneous ones due to their lower formation temperature,

and therefore they have a high strengthening potential in spite of their low volume

fraction [41].


71

6.2.2. Influence of hot-deformation on initial microstructure

The most important role of the hot-deformation is the homogenisation of the material.

Hot-deformation introduces more nucleation sites for the carbides, leading to an even

distribution of the particles during the precipitation. In addition, hot work refines the prior

austenite grain size, providing higher toughness [6].

For the sample without hot deformation (12Cr4CoWTa-780 NHD), an average PAGS of

139 ± 9 μm was measured. This value is considerably larger than the PAGS measured for

both hot-deformed samples: 33 ± 3 μm (12Cr4CoWTa-780 HD) and 32 ± 3 μm

(12Cr4CoWTa-680 HD). In Fig. 6-2 a comparison of the initial microstructure of samples

12Cr4CoWTa-780 HD and 12Cr4CoWTa-780 NHD is shown. Alloy 12Cr4CoWTa-780

HD (Fig. 6-2a) shows a uniform distribution of precipitates, which are located mainly at

prior austenite grain boundaries (Laves phase) and at sub-grain boundaries (MX particles)

and present uniform particle sizes. For the sample without hot-deformation

(12Cr4CoWTa-780 NHD) a non-uniform size distribution of precipitates is observed,

with small precipitates and large precipitates of diameters about 0.7 μm distributed

throughout the microstructure (Fig. 6-2b).

Fig. 6-2: STEM-HAADF images of the initial microstructure of sample 12Cr4CoWTa-780 HD (a) and sample

12Cr4CoWTa-780 NHD (b). The microstructure of the hot-deformed sample shows a uniform distribution of

precipitates compared to the non hot-deformed case.


72

6.2.3. Influence of tempering temperature on initial microstructure

The initial microstructures of sample 12Cr4CoWTa treated at two different tempering

temperatures are compared regarding sub-grain size dislocation density and hardness (see

Tab. 5-8). Precipitation of carbides and carbonitrides occurs during tempering. In

addition, tempering induces the recovery of the martensitic structure with transformation

into a ferritic structure, as well as the recovery of the dislocation cell structure resulting in

a sub-grain structure with a reduction of the dislocation density [ 138 ]. At higher

tempering temperatures, the development of sub-grain structures is expected together

with a reduction of the dislocation density. This is in agreement with the measurements,

the sub-grain size of sample 12Cr4CoWTa-680 HD was 380 ± 30 nm and the dislocation

density was 27.3 x 1013 m-2, whereas for sample 12Cr4CoWTa-780 HD an increment of

the sub-grain size (420 ± 32 nm) and a reduction of the dislocation density (26.2 x 1013

m-2) were measured.

For the sample without hot deformation, largest values of sub-grain size (450 ± 47 nm) as

well as smallest values of dislocation density (21.2 x 1013 m-2) were determined for the

initial state. It was observed that the hot-deformation process does not extensively affect

the size of the sub-grains (compare 12Cr4CoWTa-780 HD and 12Cr4CoWTa-780 NHD,

~ 420 nm and ~ 450 nm respectively). As described in previous works [132,138], the sub-

grain formation is related to the reduction of the dislocation density created during the

martensitic transformation and the tempering step. In our investigations, the hot-

deformation step influences the precipitate distribution (as described before) and affects

the sub-grain size formation to a lesser extent.

The hardness is closely related to the dislocation density and sub-grain size. As was

mentioned above, for higher tempering temperatures a decrease in dislocation density and

a decrease of hardness measurements were observed (Tab. 5-8). However, a slight

reduction of dislocation density and hardness was observed by increasing the tempering

temperature from 680 °C to 780°C. The high Co content of the alloy decreases the Ac1

temperature below the tempering temperature (780°C). An Ac1 temperature of 740°C was


73

calculated for this alloy. This means that retransformation to austenite took place during

tempering at 780°C leading to some martensite formation after tempering. Therefore the

presence of martensite may have contributed to the high dislocation density and hardness

for the sample tempered at 780°C.

6.2.4. Influence of hot deformation on creep strength

The creep behaviour and microstructure characteristics of samples 12Cr4CoWTa-780

NHD and 12Cr4CoWTa-780 HD differ in several aspects. Sample 12Cr4CoWTa-780 HD

presents a rupture time more than twice the rupture time of the non hot-deformed sample

(12Cr4CoWTa-780 NHD) and a reduced rupture strain. The abrupt transition from

primary to tertiary creep for the samples 12Cr4CoWTa-780 HD shown in Fig. 5-23a

could be related to degradation of the microstructure (e.g. coarsening of sub-grains and

precipitates).

As reported in Ref. [27], the dominant process of strain generation in this kind of

materials is crystallographic slip by glide of free dislocations. This means that the strain

is directly related to the dislocation evolution, as well as the sub-grain coarsening.

Moreover as Kimura reported [116], the coarsening and dissolution of fine precipitates

sometimes takes place preferentially in the vicinity of grain boundaries during creep,

which promotes the formation of localised weak zones and promotes localised creep

deformation near grain boundaries.

The investigation of the microstructure features was carried out taking into consideration

the initial stage and the final creep stage before rupture. As previously mentioned, the

creep tested samples were taken from a region about 15 mm from the fracture zone in

order to avoid the influence of the necking. In sample 12Cr4CoWTa-780 HD the

dislocation density was almost twice the value obtained for sample 12Cr4CoWTa-780

NHD (Tab. 5-9). The dislocation density and sub-grain size are directly correlated with

the hardness, being larger for the hot-deformed sample. Sample 12Cr4CoWTa-780 HD

presented a lower minimum creep rate at smaller strain in comparison to sample

12Cr4CoWTa-780 NHD.


74

6.2.5. Influence of tempering temperature on creep strength

Samples 12Cr4CoWTa-680 HD and 12Cr4CoWTa-780 HD showed large sub-grains with

low dislocation density in the interior. The sub-grains suffered a polygonalisation as a

result of the extensive deformation whereas the dislocation density and the size of the

precipitates after creep were very similar.

The creep curves for samples 12Cr4CoWTa-680 HD and 12Cr4CoWTa-780 HD showed

very similar creep behaviour. The difference lies in the steeper creep rate increase of

sample 12Cr4CoWTa-680 HD, which leads to a shorter rupture time. Sample

12Cr4CoWTa-680 HD presented a reduced ductility compared to sample 12Cr4CoWTa-

780 HD (strain 17.4%). Remarkable differences were observed in the sub-grain size

evolution during creep. For example, in 12Cr4CoWTa-680 HD the sub-grain size

increased from 380 ± 30 nm (initial stage) to 670 ± 72 nm (after creep). For sample

12Cr4CoWTa-780 HD, the sub-grain size increased from 420 ± 32 nm to 900 ± 86 nm.

Despite the lower dislocation density in 12Cr4CoWTa-680 HD compared to

12Cr4CoWTa-780 HD, the hardness measurements showed higher values for

12Cr4CoWTa-680 HD than for 12Cr4CoWTa-780 HD (Tab. 5-10). This may be related

to the fact that the sub-grain size was considerably smaller in 12Cr4CoWTa-680 HD than

in 12Cr4CoWTa-780 HD.

In Fig. 5-26 it can be observed that 12Cr4CoWTa-680 HD presented higher creep

strengths than 12Cr4CoWTa-780 HD only after short-term creep with higher stresses. A

similar behaviour was observed in [139], where a comparison of creep rupture strength of

two tempering treatments at 750°C and 800°C in 12Cr–1Mo–1W–VNb steel is shown.

A. Iseda et al. [139] showed that, for short periods of creep, the steel with low

temperature tempering shows higher creep rupture strength than the steel produced with

the high temperature tempering. The stress vs. time to rupture curves crossed over for

higher periods of time (~ 6,000 h at 650°C). This behaviour was associated to the fact

that excess dislocations accelerate recovery and recrystallisation during creep with the


75

application of stress. The authors observed that the dislocation density after tempering at

higher temperatures was too low to promote recovery and recrystallisation during creep.

In this work the quantification of dislocation density showed a rapid reduction for the

sample tempered at lower temperatures (compare Tab. 5-8 and Tab. 5-10). A mechanism

similar to that in [139] is believed to be active in the present case.

6.2.6. Conclusions for the studied 12% Cr steels

The results of the microstructure analysis of designed alloys (focussed on the precipitates

quantification) were compared with the macro-mechanical properties (creep tests 100

MPa / 650°C / 8,000 h) in order to investigate the influence of the different precipitates

on creep strength of the designed alloys, especially the influence of the M23C6 carbide.

In Section 5.2., alloy 12Cr4CoWTa was selected in order to investigate the influence of

processing parameters (hot deformation and tempering temperature) on the

microstructure evolution and creep strength. STEM investigations were focused on the

quantitative determination of dislocation density and the sub-grain size evolution by

STEM-HAADF during creep at 650°C in the early stages of creep (4,000 h), and their

correlation with the creep test results.

The conclusions of this chapter can be summarised as follows:

• ThermoCalc calculations showed to be a reliable tool for alloy development of

heat resistant steels, regarding the influence of Co and W on the Laves phase

formation and austenite stabilisation.

• Investigations of the microstructure at different creep conditions show good

agreement with the predicted phases of the thermodynamic modeling, except for

the Z-phase and Laves phase which precipitate under creep condition.


76

• Laves phases are found in the initial microstructure of alloy 12Cr4CoWTa-780

HD. No Z-phase was observed, probably due to insufficient creep times (rupture

before 4,000 h).

• No Laves phase was present in the microstructure of alloy 12Cr2CoWV-780 HD

after tempering. The Laves phase precipitates during creep condition (already

precipitates after 2,000 h creep in alloy 12Cr2CoWV-780 HD).

• Very few Z-phase precipitates after 6,150 h were observed in alloy 12Cr2CoWV-

780 HD, probably nucleating at existing (V,Ta)(C,N) to form a Cr(V,Ta)N

precipitate.

• Alloys 12Cr4CoWTa-780 HD and 12Cr2CoWV-780 HD show growth and

coarsening of Laves Phase, whereas the MX carbonitrides present very slow

growth and coarsening rates.

• Alloy 12Cr2CoWV-780 HD presents better creep properties than alloy

12Cr4CoWTa-780 HD. This difference in performance may be attributed to the

high volume fraction of the M23C6 in alloy 12Cr2CoWV in combination with the

presence of boron, providing further strengthening to the alloy.

• The hot-deformation process does not considerably affect the sub-grain size (hot-

deformed alloy 12Cr4CoWTa-780 HD = 450 ± 47 nm; non hot-deformed alloy

12Cr4CoWTa-780 NHD = 420 ± 32 nm), hence the sub-grain formation is related

primarily to the martensitic transformation and to the tempering temperature.

• Regarding alloy 12Cr4CoWTa, hot-deformed samples presented a higher density

of dislocations compared to the non hot-deformed samples. In sample

12Cr4CoWTa-780 HD, a dislocation density of 26.2 x 1013 m-2 was measured,

higher than the 21.2 x 1013 m-2 measured for sample 12Cr4CoWTa-780 NHD.


77

• The relatively high dislocation density of alloy 12Cr4CoWTa-780 HD after

tempering at 780°C could be a result of retransformation of martensite to austenite,

due to the fact that the Ac1 of this alloy (740°C) is below the temperature of the

tempering treatment. The dislocation density was 26.2 x 1013 m-2 in sample

12Cr4CoWTa-780 HD and 27.3 x 1013 m-2 in sample 12Cr4CoWTa-680 HD.

• All creep tested samples showed a significant increase of sub-grain size, as well

as a reduction of the dislocation density. As an example, the dislocation density of

the hot-deformed sample (12Cr4CoWTa-780 HD) decreased from 26.2 x 1013 m-2

(initial microstructure) to 4.7 x 1013 m-2 (microstructure after creep). Meanwhile,

the sub-grain size increased from 420 ± 32 nm (initial microstructure) to 700 ± 88

nm after creep.

• The higher creep strength of the hot-deformed sample (12Cr4CoWTa-780 HD)

compared to the non hot-deformed sample (12Cr4CoWTa-780 NHD) is related to

a more uniform particle size and distribution, as observed in STEM micrographs.

This distribution stabilises the free dislocations in the matrix and sub-grain

structure, enhancing dislocation hardening and sub-boundary hardening, which is

related to the lower minimum creep rate and higher rupture times.

• Sample 12Cr4CoWTa-680 HD presented higher creep strengths for shorter

periods of time (less than 1,000 h at 650°C) than 12Cr4CoWTa-780 HD, whereas

the stress vs. times to rupture for longer creep at lower stresses become alike.


78

7. Results of 9% Cr steels

In the present chapter four 9% Cr alloy were designed based on a combination of physical

metallurgy principles and ThermoCalc modeling. The novelty of this alloy compared

with previous works is the reduction of the W and Co content and the achievement of a

good balance of the carbide and carbonitride former elements (e.g. Nb, V and Ti) as well

as a good balance of B and N to avoid the formation of large borides.

In two alloys (9CrTi-H and 9CrTi-L), small amounts of Ti were added in order to

investigate the potential of the Ti-containing precipitates for strengthening 9% Cr

martensitic/ferritic steels for different carbon contents (between 0.05 and 0.1%).

Additionally, two alloys (9Cr-H and 9Cr-L) were designed without Ti in order to

investigate the influence of C content (between 0.05 and 0.1%) on the formation of the

different phases (e.g. carbides, carbonitrides, and Laves phase) and the strengthening of

the alloy.

7.1. Thermodynamic calculations of 9% Cr steels

The influence of Ti addition and carbon content was modeled by using ThermoCalc. In

particular the phase fields at the austenisation, tempering (780°C) and creep temperatures

(650°C) are of interest.

Based on this information the production of the samples was carried out. The Ti

containing alloys are referred as to 9CrTi. The low and high carbon content is indicated

by L = low and H = high.


79

9CrTi-H and 9CrTi-L design

Alloys 9CrTi-H and 9CrTi-L were designed to obtain ferrite, Ti-MX and Nb-MX

particles, M23C6 carbides and Laves phase at a temperature of 650°C.

Ti addition combined with C and N promotes the precipitation of Ti-MX particles. Ti-

MX particles have an extremely high stability and thus a high potential for strengthening

of martensitic/ferritic steels. As is shown in the phase diagram (Fig. 7-1), Ti-MX

precipitates are more stable compared to Z-phase, they suppress the precipitation of Z-

phase above 650°C, e.g. addition of 0.03% Ti is enough to decrease the precipitation

temperature of Z-phase below 650°C for the range of compositions investigated.

Fig. 7-1: ThermoCalc phase diagram for alloys 9CrTi-H and 9CrTi-L (F=ferrite and A= austenite,

ThermoCalc TCFe6). The austenisation temperature and the tempering temperature are indicated in the phase

diagram by black circles for each alloy. Ti-MX denotes the Ti-rich phase which contains N, C and few Nb,

whereas Nb-MX are Nb-rich particles with C and N and also few amounts of Ti and Cr.


80

ThermoCalc calculations show that alloys 9CrTi-H and 9CrTi-L contain ferrite, Ti-MX

and Nb-MX particles, M23C6 carbides and Laves phase at 650°C. Ti-MX denotes a Ti-

rich phase which contains N, C and some Nb. Ti-MX particles are more stable compared

to Z-phase at 650°C.

The phase diagram in Fig. 7-1 indicates that the precipitation of Ti-MX particles start

already in the liquid even for small additions of Ti (around 1,500°C) and there is nearly

no change in the amount of the phase with decreasing temperature. Therefore it seemed to

be difficult to control the size and distribution of the Ti-MX particles in the

microstructure. Inspite of this, the austenisation temperatures were fixed at about 50°C

above the precipitation temperature of the phase field containing Nb-MX particles for

both alloys (Tab. 7-1) avoiding regions where δ-ferrite is a stable phase at higher

temperatures.

Tab. 7-1: Austenisation temperatures from ThermoCalc.

Alloy 9CrTi-H 9CrTi-L 9Cr-H 9Cr-L

Austenisation T°C 1120 1080 1120 1080

The tempering temperature (780°C) was chosen to ductilise the hard and brittle

martensite transformed during the air-cooling after tempering. The phase diagram shows

phase fields with ferrite, Ti-MX, Nb-MX precipitates and M23C6 carbides in both alloys.

Nb-MX denotes Nb-rich particles with C and N and also few amounts of Ti and Cr. No

Laves phase is expected in the initial microstructure of both alloys.

The main difference between 9CrTi-H and 9CrTi-L is the amount of C added (0.1% and

0.05% respectively) which influences the volume fraction of M23C6 precipitates and

Laves phase. In Tab. 7-2 the volume fractions of all precipitates in both alloys at the

tempering temperature (780°C) and at the creep temperature (650°C) are shown. At

650°C alloy 9CrTi-H contains 2.03 vol.% of M23C6 and 0.62 vol.% of Laves phase,

whereas 9CrTi-L contains 0.82 vol.% of M23C6 and 0.88 vol.% of Laves phase.


81

Tab. 7-2: Volume fractions of precipitates calculated with ThermoCalc for alloys 9CrTi-H and 9CrTi-L at 780°C

and 650°C.

9CrTi-H 9CrTi-L

T°C Phases Volume Fraction (%) Volume Fraction (%)

Ferrite 97.93 99.12

M23C6 1.98 0.78

Ti-MX 0.06 0.06 780

Nb-MX 0.03 0.04

Ferrite 97.26 98.21

M23C6 2.03 0.82

Ti-MX 0.06 0.06

Nb-MX 0.03 0.03

650

Laves phase 0.62 0.88

• 9Cr-H and 9Cr-L design

The chemical composition of alloys 9Cr-H and 9Cr-L were designed to have ferrite, V-

MX, Nb-MX precipitates, M23C6 carbides and Laves phase at 650°C. ThermoCalc

calculations show (see Fig. 7-2) that alloys 9Cr-H and 9Cr-L indeed contain ferrite, V-

MX, Nb-MX precipitates, M23C6 carbides and Laves phase at 650°C. V-MX corresponds

to V-rich precipitates which contain Nb, N and C and few Fe and Cr, whereas Nb-MX

refers to Nb-rich particles with C, Cr and N and also few amounts of V


82

Fig. 7-2: ThermoCalc phase diagram for alloys 9Cr-H and 9Cr-L (F=ferrite and A= austenite, ThermoCalc

TCFe6). The austenisation temperature and the tempering temperature are indicated in the phase diagram by

black circles for each alloy. V-MX is V-rich phase containing Nb, N and C and few Fe and Cr, whereas Nb-

MX denotes Nb-rich particles with C, Cr and N and also few amounts of V.

The phase diagram in Fig. 7-2 shows that Z-phase is more stable compared to V-MX

precipitates at 650°C. In Ref [19] was demonstrated that V-MX precipitates are gradually

transformed into Z-phase, which leads to an early consumption of the V-MX particles in

the region adjacent to the prior austenite grain boundaries, decreasing the creep strength.

Despite the calculations, several studies had reported that 9% Cr steels do not suffer from

abundant formation of Z-phase after long-term creep [88,140] due to the slow kinetics of

the Z-phase precipitation.


83

The austenisation temperatures were chosen about 50°C above the phase field containing

Nb-MX particles (1120°C and 1080°C for the 9Cr-H and 9Cr-L alloy, respectively), in

order to obtain a fully austenitic field and ensure a completely martensitic transformation

(avoiding δ-ferrite see Fig. 7-2).

The tempering temperature (780°C) was chosen to temper the martensite transformed

during air-cooling after the austenisation treatment. The stable phases present in both

alloys were ferrite, M23C6 carbides, V-MX and Nb-MX precipitates. No Laves phase is

expected in the initial microstructure of the 9Cr alloys.

The main difference between 9Cr-H and 9Cr-L lies in the amount of C added (0.1% and

0.05% respectively) which influences the volume fraction of M23C6 precipitates and

Laves phase. The volume fractions for all precipitates were calculated at the tempering

temperature 780°C and at the creep temperature 650°C (Tab. 7-3). At 650°C alloy 9Cr-H

contains 2.11 vol.% M23C6 and 0.67 vol.% Laves phase, whereas 9Cr-L contains 0.96

vol.% M23C6 and 0.93 vol.% Laves phase.

Tab. 7-3: Volume fractions of precipitates calculated with ThermoCalc for alloys 9Cr-H and 9Cr-L at 780°C and

650°C.

9Cr-H 9Cr-L

T°C Phases Volume Fraction (%) Volume Fraction (%)

Ferrite 97.89 99.02

M23C6 2.05 0.92

V-MX 0.03 0.02 780

Nb-MX 0.03 0.04

Ferrite 97.14 98.02

M23C6 2.11 0.96

Z-phase 0.06 0.06

Nb-MX 0.02 0.03

650

Laves phase 0.67 0.93


84

7.2. Alloy production

The four designed alloys (9CrTi-H, 9CrTi-L, 9Cr-H and 9Cr-L) were prepared by

vacuum induction melting with masses of about 4 kg.

The analysed chemical compositions of the alloys are shown in Tab. 7-4.

The samples were produced with the following parameters:

• Hot-rolling at 1150°C with subsequent air cooling (66% final deformation).

• Austenisation heat treatment for 0.5 h followed by air-cooling (the austenisation

temperatures were calculated with ThermoCalc)

- Alloys 9CrTi-H and 9Cr-H austenitised at 1120°C

- Alloys 9CrTi-L and 9Cr-L austenitised at 1080°C

• Tempering for 2 h at 780°C with subsequent air-cooling

Fig 7-3 shows a scheme of the production of the 9% Cr steels

Fig. 7-3: 9% Cr steels heat treatments scheme.


85

Tab. 7-4: Analysed chemical composition of the produced alloys (wt%).

Alloys B C Co Cr Mn N Nb Si Ti V W

9CrTi-H 0.007 ± 0.005

0.106 ± 0.01

1.01 ± 0.10

9.08 ± 0.50

0.53 ± 0.10

0.008 ± 0.002

0.030 ± 0.01

0.36 ± 0.05

0.030 ± 0.01

0.15 ± 0.02

1.93 ± 0.10

9CrTi-L 0.008 ± 0.005

0.047 ± 0.01

1.01 ± 0.10

8.90 ± 0.50

0.53 ± 0.10

0.007 ± 0.002

0.030 ± 0.01

0.36 ± 0.05

0.035 ± 0.01

0.15 ± 0.02

1.90 ± 0.10

9Cr-H 0.009 ± 0.005

0.108 ± 0.01

0.99 ± 0.10

8.81 ± 0.50

0.49 ± 0.10

0.005 ± 0.002

0.030 ± 0.01

0.38 ± 0.05 - 0.15 ±

0.02 2.00 ± 0.10

9Cr-L 0.013 ± 0.005

0.051 ± 0.01

1.01 ± 0.10

8.97 ± 0.50

0.48 ± 0.01

0.005 ± 0.002

0.035 ± 0.01

0.37 ± 0.05 - 0.15 ±

0.02 1.98 ± 0.10


86

7.3. Microstructure evolution (precipitates and sub-grain size)

Light microscopy of the investigated alloys after heat treatment shows a

martensitic/ferritic matrix and precipitates. STEM investigations were carried out to

quantify the microstructure features in the initial stage and after creep (to rupture).

• Alloys 9CrTi-H and 9CrTi-L (Influence of Ti)

Initial microstructure: In the initial condition (after tempering at 780°C/ 2h) both alloys

presented a martensitic/ferritic matrix with high density of interfaces, such as prior

austenite grain boundaries, prior lath boundaries and sub-grain boundaries, as well as

high dislocation density.

The average sub-grain size was measured for both alloys (Tab 7-5). For alloy 9CrTi-H

the sub-grain was 401 ± 44 nm, slightly smaller than the sub-grain size of 9CrTi-L with

an average size of 433 ± 25 nm.

Tab. 7-5: Sub-grain size and hardness of alloys 9CrTi-H and 9CrTi-L at initial stage.

Alloy Sub-grain size(nm) HV10

9CrTi-H initial 401 ± 44 246 ± 2

9CrTi-L initial 433 ± 25 223 ± 1

The initial microstructure in both alloys showed M23C6 carbides and Nb-rich and Ti-rich

precipitates. The particles were identified by a combination of DF and EDS (since the

identification through chemical analysis of C and N could not precisely be obtained in

this study, they are generally described as carbonitrides).

The alloy 9CrTi-H (Fig. 7-4a) shows a high amount of M23C6 precipitates with an

average size of 79 ± 5 nm (see Tab. 7-6). Fig. 7-4b shows the M23C6 carbides for alloy

9CrTi-L with an average size of about 89 ± 4 nm (Tab. 7-5). The quantity of M23C6

carbides presented in 9CrTi-L was much smaller than that of the precipitates in 9CrTi-H.


87

Few large W and Fe rich particles (about 500 nm) were found in alloy 9CrTi-L (Fig. 7-

4b).

Fig. 7-4: STEM-HAADF micrographs of initial microstructure of alloy 9CrTi-H (A) and alloy 9CrTi-L (B).

White arrows point at the M23C6 precipitates in both micrographs. Alloy 9CrTi-L shows large particles rich in

W and Fe (possibly FeW2B).

Tab. 7-6: Average size of precipitates in alloys 9CrTi-H and 9CrTi-L at initial stage (time in hours and size in

nanometers).

9CrTi-H 0 h 9CrTi-L 0 h M23C6 79 ± 5 M23C6 89 ± 4 Ti-MX 30 ± 1 Ti-MX 57 ± 6 Nb-MX 29 ± 2 Nb-MX 27 ± 2

Laves phase 0 Laves phase 0

Nano-sized Nb-MX particles with spheroidal shape and Ti-MX precipitates with

rhomboidal shape were detected in both alloys (Fig 7-5 and Fig. 7-6). TEM-EDS

measurements showed that the Nb-MX precipitates are Nb-rich particles which contain C

and N and few amounts of Ti and Cr, whereas Ti-MX are Ti-rich precipitates which

contain N, C, Nb and Cr. Both Nb-MX and Ti-MX are generally described as

carbonitrides. The average size of Nb-MX particles was very similar in both investigated

alloys (Tab. 7-6).


88

In alloy 9CrTi-H the Nb-MX particle size was about 29 ± 2 nm, whereas in alloy 9CrTi-L

the mean particle size was 27 ± 2 nm. The identified Ti-MX particles in alloy 9CrTi-H

(Fig. 7-6) showed a particle size of 30 ± 1 nm, whereas the average particle size of Ti-

MX precipitates in alloy 9CrTi-L was about 57 ± 6 nm.

Fig. 7-5: STEM-HAADF micrograph of the initial microstructure of alloy 9CrTi-H (A) (nano-sized Nb-MX

precipitates are indicated by white arrows; black arrows point at the Ti-MX particles). EDS spectrum of the

encircled Nb-MX precipitate (B).

Fig. 7-6: STEM-HAADF micrograph of initial microstructure of alloy 9CrTi-L (A) showing Ti-rich MX

precipitates (white arrows), the M23C6 precipitates are pointed at by black arrows and EDS spectrum of the

encircled Ti-MX particle (B).


89

Few large TiN precipitates of about 700 nm were observed in 9CrTi-L (Fig. 7-7), no large

Ti particles were found in 9CrTi-H.

The average hardness for both alloys is shown in Tab. 7-5. The mean hardness value for

alloy 9CrTi-H was about 246 ± 2 HV10, whereas the hardness measured for alloy 9CrTi-L

was 223 ± 1 HV10.

Fig. 7-7: STEM-HAADF micrograph initial microstructure of alloy 9CrTi-L (A) showing a large Ti-rich

precipitate and the M23C6 precipitates (white arrows) and EDS spectrum of the Ti-rich particle (B).

Microstructure after creep: The sub-grain size measured for alloy 9CrTi-H after 7,253

h creep at 650°C / 101 MPa was about 821 ± 69 nm (Tab. 7-7). For alloy 9CrTi-L the

sub-grain size was even larger, reaching 1011 ± 94 nm after 2,154 h creep at 650°C/101

MPa.

Tab. 7-7: Sub-grain size and hardness of alloys 9CrTi-H and 9CrTi-L after creep at 650°C.

Samples Stress (MPa) Rupture time (h) Sub-grain size (nm) HV10

9CrTi-H 101 7,253 821 ± 69 224 ± 2

9CrTi-L 101 2,154 1011 ± 94 194 ± 3


90

The hardness values were consistent with the enlargement of the sub-grain size for both

alloys. A decrease from 246 ± 2 HV10 to 224 ± 2 HV10 was observed for alloy 9CrTi-H,

whereas a decrease from 223 ± 1 HV10 to 194 ± 3 HV10 was measured for alloy 9CrTi-L.

Both alloys contained M23C6 carbides, Ti-MX and Nb-MX particles and Laves phase

(Tab. 7-8). M23C6 precipitates were still the most abundant particles. The average particle

size of M23C6 precipitates (Fig. 7-8, white arrows) was 97 ± 5 nm for alloy 9CrTi-H (Fig.

7-8a) and 106 ± 8 nm for alloy 9CrTi-L (Fig. 7-8b). The lath or block boundaries were

often pinned by M23C6 carbides.

Tab. 7-8: Average size of precipitates in alloys 9CrTi-H and 9CrTi-L (time in hours and size in nanometers) after

creep (650°C / 101MPa).

9CrTi-H 7,253 h 9CrTi-L 2,154 h M23C6 97 ± 5 M23C6 106 ± 8 Ti-MX 34 ± 2 Ti-MX 65 ± 4 Nb-MX 33 ± 3 Nb-MX 28 ± 2

Laves phase 411 ± 32 Laves phase 347 ± 40

Fig. 7-8: STEM-HAADF micrograph of alloy 9CrTi-H (A) after creep (7,253h / 101MPa / 650°C) and STEM-

HAADF micrograph of alloy 9CrTi-L (B) after creep (2,154h / 101MPa / 650°C) with M23C6 precipitates

(white arrows) and Laves phase (black arrows).


91

Larger Laves phase particles (black arrows) were detected near the M23C6 carbides, as

shown in Fig. 7-8 on prior austenite grain boundaries and lath or block boundaries. Laves

phase formed and grew under creep condition after several hundred hours. Fig. 7-9 shows

an example of the particle identification procedure. The main elements detected by TEM-

EDS in the Laves phase were W, Fe and some Cr. The average particle size of Laves

phase for alloy 9CrTi-H was 411 ± 32 nm and 347 ± 40 nm for alloy 9CrTi-L (Tab 7-8).

Fig. 7-9: STEM-HAADF micrograph of alloy 9CrTi-L (A) after creep (2,154h / 101 MPa / 650°C), diffraction

pattern of the encircled Laves phase particle (B), and EDS spectrum of Laves phase particle (C).

The average particle size of Ti-MX precipitates in alloy 9CrTi-H after creep was 34 ± 2

(initially 30nm, 13% growth) nm. For alloy 9CrTi-L the Ti-MX particles size was 65 ± 4

nm (initially 57 nm, 14% growth).

Fig. 7-10 shows the tensile creep curves of alloys 9CrTi-H and 9CrTi-L at 101 MPa /

650°C. Alloy 9CrTi-H shows a decrease in the minimum creep rate and increase of the

time to rupture compared to 9CrTi-L.


92

Fig. 7-10: Tensile creep curves comparing the creep strength of alloys 9CrTi-H and 9CrTi-L at 101 MPa /

650°C, (A) strain vs. time, (B) creep vs. strain.

• Alloys 9Cr-H and 9Cr-L (Influence of C)

Initial microstructure: A martensitic/ferritic matrix with a high density of interfaces and

a high dislocation density was observed in the initial microstructure in both alloys. Sub-

grain size and hardness values are shown in Tab. 7-9. The sub-grain size for alloy 9Cr-H

was 426 ± 27 nm and the hardness value was about 248 ± 3 HV10. For alloy 9Cr-L the

sub-grain size was 403 ± 33 nm and the hardness was 221 ± 2 HV10.

Both alloys showed an initial microstructure with Nb-MX, V-MX precipitates and M23C6

carbides. Alloy 9CrTi-H (Fig. 7-11a) showed a higher amount of M23C6 precipitates with

an average size of 78 ± 3 nm, whereas for alloy 9CrTi-L (Fig. 7-11b) the average size of

the M23C6 carbides was about 97 ± 4 nm (Tab. 7-10).

Tab. 7-9: Sub-grain size and hardness of alloys 9Cr-H and 9Cr-L at initial stage.

Samples Sub-grain size (nm) HV10

9Cr-H initial 426 ± 27 248 ± 3

9Cr-L initial 403 ± 33 221 ± 2


93

Fig. 7-11: STEM-HAADF micrographs of initial microstructure of alloy 9Cr-H (A) and alloy 9Cr-L (B) with

M23C6 precipitates (white arrows).

Tab. 7-10: Average size of precipitates in alloys 9Cr-H and 9Cr-L at the initial stage (time in hours and size in

nanometers).

9Cr-H 0 h 9Cr-L 0 h M23C6 78 ± 3 M23C6 97 ± 4 V-MX 30 ± 1 V-MX 30 ± 2

Nb-MX 29 ± 2 Nb-MX 25 ± 2 Laves phase 0 Laves phase 0

Nano-sized Nb-MX particles with a spheroidal shape and V-MX precipitates with a plate-

like shape were identified in both alloys (see Fig. 7-12 and Fig. 7-13). TEM-EDS

measurements showed that the Nb-MX are Nb-rich particles which contain C, N, V and

Cr. V-MX are V-rich precipitates which contain N, C, Nb and Cr. Both Nb-MX and V-

MX are generally described as carbonitrides.

The mean particle size of the Nb-MX precipitates in alloy 9Cr-H was 29 ± 2 nm, whereas

in alloy 9Cr-L (Fig. 7-12) the Nb-MX particle size was 25 ± 2 nm (Tab. 7-10). A similar

particle size as for Nb-MX particles was measured for the plate-like V-MX particles whit

an average diameter of 30 ± 1 nm for alloy 9Cr-H and 30 ± 2 for alloy 9Cr-H. (Fig. 7-13).

For the plate-like V-MX particles two perpendicular axes were measured (a and b) and an

average diameter d = (a+b)/2 was calculated (see section 4.4.).


94

Fig. 7-12: STEM-HAADF micrograph of the initial microstructure of alloy 9Cr-L (A) with Nb- MX particles

and EDS spectrum of the encircled Nb-MX particle (B).

Fig. 7-13: STEM-HAADF micrographs of the initial microstructure of alloy 9Cr-L (A) with V-MX particles

(white arrows) and Nb-MX particles (black arrows) and EDS spectrum of the encircled V-MX particle (B).

Microstructure after creep: The sub-grain size of the 9Cr-H alloy increased from 426 ±

27 nm to 689 ± 68 nm after 7,987 h creep at 650°C / 101 MPa (Tab. 7-11). For alloy 9Cr-

L the increase in sub-grain size was slightly smaller compared to 9Cr-H, ranging from

403 ± 33 nm in the initial state to 647 ± 46 nm after 10,168 h creep at 650°C / 125 MPa

(Tab. 7-11). In the initial state alloy 9Cr-H presented a mean hardness value of 248 ± 3

HV10; after creep the hardness was about 244 ± 2 HV10. For alloy 9Cr-L the initial

hardness was 221 ± 2 HV10, whereas after creep (to rupture) the hardness was 220 ± 3

HV10.


95

Tab. 7-11: Sub-grain size and hardness in alloys 9Cr-H and 9Cr-L after creep at 650°C.

Samples Stress (MPa) Rupture time (h) Sub-grain size (nm) HV10

9Cr-H 101 7,987 689 ± 68 244 ± 2

9Cr-L 125 10,168 647 ± 46 220 ± 3

Both creep-tested alloys contained M23C6 carbides, V-MX and Nb-MX particles and

Laves phase (Tab. 7-12). The M23C6 precipitates showed an average particle size of 103 ±

6 nm for alloy 9Cr-H (Fig. 7-14a) and of 112 ± 8 nm for alloy 9Cr-L (Fig. 7-14b). In both

cases the M23C6 carbides were mostly placed on the prior austenite boundaries and on

lath or block boundaries. Fig. 7-15 shows an example of the M23C6 carbide identification.

The EDS measurement showed Cr, W and Fe as main components of this carbide.

Tab.7-12: Average size of precipitates from alloys 9CrTi-H and 9CrTi-L (time in hours and size in nanometers)

under creep condition (9Cr-H 650°C / 101MPa and 9Cr-L 650°C / 125MPa).

9Cr-H 7,987 h 9Cr-L 10,168 h M23C6 103 ± 6 M23C6 112 ± 8 V-MX 31 ± 1 V-MX 31 ± 2

Nb-MX 31 ± 1 Nb-MX 29 ± 3 Laves 379 ± 48 Laves 354 ± 42


96

Fig. 7-14: STEM-HAADF micrograph of alloy 9Cr-H (A) after creep (7,987h / 101MPa / 650°C) and STEM-

HAADF micrograph of alloy 9Cr-L (B) after creep (10,168h / 125MPa / 650°C) white arrows indicate M23C6

precipitates; black arrows indicate Laves phase.

Fig. 7-15: (A) STEM-HAADF micrograph of alloy 9Cr-L after creep (10,168h/ 125 MPa / 650°C). (B)

Diffraction pattern of the encircled M23C6 particle. (C) EDS spectrum of the encircled M23C6 particle.

Fig. 7-16 shows the interaction of dislocations and sub-grains with the precipitates,

especially with M23C6 carbides (sample 9Cr-L after 10,168 h creep at 650°C/125MPa).

The average particle size of Laves phase for the 9Cr-H alloy after creep was 379 ± 48 nm

and 347 ± 40 nm for alloy 9CrTi-L (Tab. 7-12).

V-MX precipitates were very stable after creep. The average particle size for both alloys

was 31 ± 2 nm (Tab. 7-12)


97

Fig. 7-16: STEM-HAADF micrograph of sample 9Cr-L after 10,168 h creep at 650°C / 125MPa (inversed

contrast). Black arrows point at Laves phase particles, white arrows indicate the M23C6 carbides. Sub-grains

and dislocations are often pinned by the M23C6 carbides.

• Creep results

Creep test results are shown in Fig. 7-17 for all investigated alloys. The present creep

tests at 650°C show the longest time to rupture for alloy 9Cr-L compared to all alloys

investigated (10,168 h at 650°C / 125MPa).

Alloys 9CrTi-H and 9Cr-H showed similar creep strengths with time to rupture about

7,500 h at 650°C / 101 MPa.


98

Alloy 9CrTi-L showed the lowest time to rupture (2,154 h at 650°C / 101MPa).

Comparing with the reference steel P92, alloy 9Cr-L showed promising results with

higher time to rupture at 650°C / 125MPa. 9CrTi-H and 9Cr-H showed similar rupture

times as P92 at 650°C / 101MPa.

Fig. 7-17: Results of the tensile creep tests at 650°C showing time to rupture as a function of applied stress for

the four investigated alloys. The alloy 9Cr-L as well as 9CrTi-H and 9Cr-H show the highest creep strength.

Corresponding data for the P92 steel [42] are shown as reference.

Discussion of 9% Cr steels

99

8. Discussion of 9% Cr steels 8.1. Microstructure evolution (precipitates and sub-grain size)

• Alloys 9CrTi-H and 9CrTi-L (Influence of Ti)

Initial microstructure: The high dislocation density is produced when martensite forms

during air-cooling after the austenisation heat treatment. During tempering, the

precipitation of solute atoms occurs, as well as recovery of the dislocation cell structure,

resulting in a sub-grain structure [101].

STEM investigations of the initial microstructure were carried out in order to measure the

sub-grain size (see Tab 7-5) because sub-grain boundary hardening is one of the most

important strengthening mechanisms for this kind of materials. The average sub-grain

size measured for alloy 9CrTi-H was slightly smaller than the sub-grain size of 9CrTi-L,

hence the sub-grain formation is related primarily to the martensitic transformation and to

the tempering temperature

The initial microstructure of both alloys showed M23C6 carbides and MX particles (Nb-

rich and Ti-rich precipitates) which are in agreement with the thermodynamic equilibrium

calculations at the tempering temperature (Fig. 7-1, 780°C).

The M23C6 carbides were the most abundant precipitates in the microstructure. The M23C6

precipitates were mostly placed on prior austenite grain boundaries and lath or block

boundaries. At such preferential sites, the effective surface energy is lower, thus

diminished the free energy barrier and facilitating nucleation [114].

The quantity of M23C6 carbides (white arrows, Fig. 7-4a and 7-4b) presented in 9CrTi-L

was much smaller than that of the precipitates in 9CrTi-H, which clearly reflects the

effect of the carbon content in this alloy. The microstructure observations are consistent

with the thermodynamic equilibrium calculations (Tab 7-2).


100

Large W and Fe rich particles (about 500 nm) were found in alloy 9CrTi-L (Fig. 7-4b).

According to Ref. [72] these particles may probably correspond to undissolved borides

(FeW2B). Such particles were not found in alloy 9CrTi-H probably due to the higher

austenisation temperature (1120°C), which allows further dissolution of borides.

Nano-sized Nb-MX particles with spheroidal shape and Ti-MX precipitates with

rhomboidal shape were detected in both alloys. They were mostly formed on sub-grain

boundaries and within the sub-grains. According to Taneike [82] the nano-sized MX

particles precipitate more homogeneously than M23C6 carbides or Laves phase due to the

small misfit between the crystallographic structure of the MX with the matrix.

The Ti-MX particles in alloy 9CrTi-L showed a larger particle size compared to alloy

9CrTi-H. The difference in the mean particle size may be related to the higher

austenisation temperature of alloy 9CrTi-H. As a consequence higher amounts of large Ti

particles are dissolved and a finer precipitation of Ti-MX carbonitrides due to higher

supersaturation during subsequent tempering is reached. Consequently no large TiN

particles were found in 9CrTi-H, whereas a few large TiN precipitates (700 nm) were

observed in 9CrTi-L.

The mean hardness value for alloy 9CrTi-H was about 246 ± 2 HV10, slightly larger than

the hardness measured for alloy 9CrTi-L (223 ± 1 HV10), which is consistent with the

carbon content variation, corresponding to secondary hardening (precipitation hardening).

Microstructure after creep: For both alloys the microstructure after creep showed large

sub-grains as a result of the extensive deformation and to the recovery of the sub-grain

structure (Tab. 7-7). However, the martensite lath structure could still be distinguished

(Fig. 7-8). In both cases the sub-grain sizes were more than twice the size of the initial

microstructure. The hardness values were consistent with the enlargement of the sub-

grain size showing a decrease in the average hardness in both cases (Tab. 7-7).


101

For both alloys the observed phases are in good agreement with the ThermoCalc

calculations at 650°C with stable phase fields containing M23C6 carbides, Ti-MX and Nb-

MX particles and Laves phase. M23C6 precipitates were still the most abundant particles

in both alloys (see Tab.7-2) and they often pinned the lath or block boundaries

It was observed (Fig. 7-8b) that the recombination of lath boundaries caused the

disappearance of some lath boundaries leaving a row of M23C6 carbides in the matrix (see

encircled area).

Laves phase formed and grew under creep conditions after several hundred hours on prior

austenite grain boundaries and lath or block boundaries. The larger size of Laves phase

particles for alloy 9CrTi-H may be related to the longer creep times (7,253 h creep at

650°C/ 101 MPa compared with 2,154 h creep at 650°C/101 MPa for sample 9CrTi-L).

Ti-MX particles in alloy 9CrTi-H showed a low coarsening rate compared to Ti-MX in

alloy 9CrTi-L after creep. An explanation for the different coarsening behaviour may be

related to a less uniform particle size of the Ti-MX precipitates measured in alloy 9CrTi-

L (Tab. 7-8, Fig. 7-6) compared to the Ti-MX particles measured in 9CrTi-H. The

smaller particles have a higher surface to volume ratio than larger particles, thus smaller

particles are less stable than larger particles of the same phase. An increase in the mean

particle size will thus reduce the total free energy of the system and this reduction in free

energy is the driving force for the coarsening reaction.

The sub-grain enlargement as well as the coarsening of the precipitates is directly

correlated to the creep strength. Under long-term conditions the precipitates coarsen with

deceasing particle number during creep deformation leading to a decreased in the pinning

of sub-grain boundaries.

Alloy 9CrTi-H shows a decrease in the minimum creep rate and increases the time to

rupture compared to 9CrTi-L (Fig. 7-10). This result suggests slow coarsening of fine

M23C6 carbides in alloy 9CrTi-H, hence a large pinning force for boundary migration is


102

maintained up to long times so that the onset of tertiary creep is retarded to longer times.

This effectively decreases the minimum creep rate and increases the time to rupture as

shown in Fig. 7-10. Crept tested samples showed a reduction of hardness, which

correlates with the observed enlargement of sub-grain size and the decrease of dislocation

density; both softening the alloy (Tab 7-5. and Tab. 7-7).

• Alloys 9Cr-H and 9Cr-L (Influence of C)

Initial microstructure: During tempering, the precipitation of solute atoms occurs, as

well as recovery of the dislocation cell structure, resulting in a sub-grain structure [118].

The initial sub-grain size for both samples is very similar, though the austenisation

temperature was 40°C higher for alloy 9Cr-H. This suggests that the austenisation

temperature has not a dominant effect on the sub-grain formation.

The observed precipitates (Nb-MX, V-MX precipitates and M23C6 carbides) are in good

agreement with the ThermoCalc result for tempering (Fig. 7-2, 780°C). The M23C6

precipitates were mostly placed on prior austenite grain boundaries and lath or block

boundaries. In agreement with the ThermoCalc calculations (Tab 7-3) alloy 9Cr-H with

high C content presented a higher volume fraction of M23C6 carbides compared to the low

C alloy (see Fig. 7-11a and 7-11b).

Nano-sized Nb-MX particles with a spheroidal shape and V-MX precipitates with a plate-

like shape were identified in both alloys. The Nb-MX and V-MX particles were mostly

located within the sub-grains or at the sub-grain boundaries pinning dislocations or sub-

grain boundaries (Fig. 7-12).

Microstructure after creep: Both alloys showed enlarged sub-grain sizes after creep

(Tab. 7-11). The hardness values in both cases remain almost constant. As mentioned

before the Vickers hardness correlates with the sub-grain size and the dislocation density

in the material. For the 9Cr alloys the increase in sub-grain size was smaller than for the


103

9CrTi alloys. This effect would suggest an effective pinning of dislocations by the

precipitates for this alloys (see Fig. 7-16).

The observed phases after creep (M23C6 carbides, V-MX and Nb-MX particles and Laves

phase) are in good agreement with the ThermoCalc calculations except for the Z-phase

phase (Tab. 7-3). The Z-phase was not detected in the microstructure after creep probably

due to the slow precipitation kinetics of the Z-phase in the 9% Cr steels, as was

previously describes by Danielsen et al. in Ref [92].

The M23C6 particles were mostly placed on the prior austenite boundaries and on lath or

block boundaries (Fig. 7-16). The particle size of M23C6 carbides in alloy 9Cr-L showed a

relatively slow coarsening rate after 10,168 h / 125 MPa compared to the M23C6 carbides

of the other alloys, the increase in the particle size was from 97 ± 4 nm at initial state

(after tempering) to 112 ± 8 nm after creep. As example the M23C6 carbides showed 19%

growth in alloy 9Cr-L, whereas in alloy 9Cr-H the M23C6 particles showed 32% growth,

from 78 ± 3 nm (after tempering) to 103 ± 4 nm after creep (7,987 h / 101 MPa).

Large Laves phase particles (see Fig. 7-14) were often observed near the M23C6 carbides

(black arrows) which are placed on prior austenite grain boundaries and lath or block

boundaries. The measurements of Laves phase after creep suggests a lower growth of

Laves phase particles in alloy 9Cr-L compared to the Laves phase growth in alloy 9Cr-H.

An explanation for this behaviour may be related to the competitive growth between

M23C6 carbides and the Laves phase. Alloy 9Cr-H showed higher amount of M23C6

carbides which nucleate and growth on prior austenite grain boundaries and lath or block

boundaries, which are the same sites where Laves phase were observed. This suggests

that M23C6 carbides nucleate and growth first on this low energetic sites and restrict the

nucleation of Laves phase. Less nucleus of Laves phase particles with a fast growth

behaviour are expected for high carbon alloys such as the 9Cr-H.

V and Nb MX particles were mostly placed on sub-grain boundaries and within the sub-

grain. V and Nb MX precipitates showed very slow coarsening rates in both alloys. This


104

observation suggests that both MX precipitates act as obstacles for the sub-grain

boundaries, thus the fine distribution of V and Nb MX particles may effectively exerts

pinning force for the migration of sub-grain boundaries up to long times during creep.

• Creep

Alloy 9Cr-L showed the best creep performance of all alloys investigated (10,168 h at

650°C / 125MPa). Alloy 9CrTi-L showed the lowest time to rupture (2,154 h at 650°C /

101MPa).

The microstructure investigations revealed that the 9Cr-L presented the lowest sub-grain

growth, very slow coarsening of MX carbonitrides and the lowest Laves phase coarsening

rate compared to all other alloys investigated.

The measured small sub-grain sizes suggest that the precipitates may provide a large

pinning force that reduced the boundary migration up to longer times, thus the tertiary

creep is retarded to longer times. This may be an explanation to the high creep strength of

alloy 9Cr-L (Fig. 7-17).

The 9CrTi-L alloy presented the largest sub-grain size growth (from 433 to 1011 nm after

only 2,154 h), large TiN particles (700 nm) and some FeW2B-like inclusions. The Ti-MX

particles showed a relatively high coarsening rate (with mean particle size growth from

57 ± 6 to 65 ± 4 nm after only 2,154 h) compared to the other alloys investigated. Both

effects reduce the effective pinning of dislocations by the precipitates for the 9CrTi-L

alloy.


105

8.2. Conclusions of studied 9% Cr steels

In the present section the microstructure before and after creep was investigated by

STEM-HAADF with respect to the evolution of the sub-grains and the precipitate

distribution for four newly designed 9% Cr heat resistant alloys. Two sets of alloys were

studied: 9Cr alloys with high (9Cr-H / 0.1%C) and low carbon contents (9Cr-L / 0.05%C)

and 9Cr alloys containing ~ 0.03Ti% and also high (9CrTi-H / 0.1%C) and low (9CrTi-L

/ 0.05%C) carbon. Correlations between the microstructure evolution and the macro-

mechanical properties were studied. The conclusions of this part of the study are

summarised as follows:

• ThermoCalc calculations showed to be a reliable tool for alloy development of heat

resistant steels. Processing parameters (austenisation and tempering temperatures)

were defined based on the phase diagram information. Investigations of the

microstructure showed good agreement with the predicted phases of the

thermodynamic modeling.

• As predicted by the thermodynamic modeling, no Laves phase precipitates were

found in the initial microstructure for the four investigated alloys.

• Ti-MX precipitates are more stable compared to Z-phase for Ti-containing alloys.

According to ThermoCalc calculations, addition of ~ 0.03% Ti effectively decreases

the precipitation temperature of Z-phase below 650°C for the investigated alloys.

• The volume fraction of precipitated M23C6 carbides is directly related to the carbon

content of the alloys. ThermoCalc calculations demonstrates that the volume fraction

of M23C6 carbides is doubled for alloys with 0.1% of carbon compared to alloys with

0.05% carbon content at 650°C.


106

• Ti-MX precipitates with rhomboidal shape were detected in 9CrTi alloys in the initial

state and after creep. Few large TiN precipitates (700 nm) were observed in the

9CrTi-L alloy. For the high carbon 9CrTi alloy no large TiN particles were observed.

• Recombination of lath boundaries was observed in all alloys after creep. This

phenomenon was more widespread in alloy 9CrTi-L.

• In 9CrTi alloys the average sub-grain size after creep was more than twice the size in

the initial microstructure. The hardness values were consistent with the increase of the

sub-grain size and experienced a decrease after creep exposure.

• For 9Cr alloys coarsening of sub-grain size was smaller than for 9CrTi alloys.

• Nano-sized Nb-MX particles with a spheroidal shape and V-MX particles with a

plate-like shape were observed in alloy 9Cr-H and alloy 9Cr-L at the initial stage and

after creep. Precipitates showed low coarsening rates and were mostly placed within

the sub-grains or at the sub-grain boundaries frequently pinning the dislocations or

sub-grains boundaries.

• Alloy 9Cr-L showed fine and stable M23C6 carbides and MX precipitates, as well as

the minimal coarsening of Laves phase after 10,168 h at 650°C / 125 MPa.

• Alloy 9Cr-L showed the highest creep strength of all the investigated alloys.

Final conclusion and perspectives

107

9. Final conclusion and perspectives

9-12% Cr heat resistant steels were designed supported by thermodynamic modeling.

ThermoCalc calculations showed to be a reliable tool for alloy development. The

influence of the alloying elements (12 components alloys) as well as the predicted

equilibrium phases were in accordance with the experimental observations. The modeling

also provided valuable information for the adjustment of the processing parameters

(austenisation and tempering temperatures).

For the 12% Cr designed alloys it was observed that microstructures presenting MX,

M23C6 and Laves phase showed high creep strength (8,000 h / 650°C / 100 MPa). The

reasons for the creep rupture were found in the formation of Z-phase and the extended

coarsening of Laves phase. The addition of boron had a large influence on the coarsening

kinetics of M23C6 improving the creep strength. Quantitative investigations of

microstructure evolution showed that in particular the initial distribution of precipitates

and the coarsening of sub-grains affect the creep rupture life of the alloy.

Following considerations for future developments -based on the observations of this work

on 12% Cr alloys- can be listed:

Laves phase: The possibility of producing a much finer precipitation of the Laves phase

as well as a reduced coarsening of this phase should be investigated. The former could be

achieved by a better control of the distribution of Cu in the alloy providing more

nucleation sites. A main challenge is to find a combination of alloying elements that

reduce the growth and the coarsening of Laves phase at 650°C.

Z-Phase: The formation of Z-Phase in 12% Cr steels at 650°C is a major challenge for the

application of such steels in power plants. Because the Z-phase is the most stable

carbonitride in 12% Cr heat resistant alloys, the possibility of precipitating the Z-phase as

a fine dispersion in the initial microstructure may be of interest. Investigations in this

direction are being carried out in Denmark by Hald et al. [141] in 12% Cr alloys with

considerably expensive master alloys containing high Ta and Co contents.


108

In order to avoid Z-phase formation the content of Cr could be reduced. For Cr contents

below 11% the Z-phase formation can be retarded considerably. The design of 11% Cr

heat resistant alloys could be a good compromise regarding creep strength and oxidation

resistance.

Alloying with Ti could avoid the formation of Z-phase at 650°C. Ti consumes nitrogen in

the formation of Ti-MX carbonitrides, which is a main element for the formation of the

Z-Phase. The use of Ti implies a further technological challenge for the austenisation of

the alloy in order to avoid the formation of large TiN precipitates in the melting

conditions.

M23C6: It was observed that the coarsening of M23C6 is higher for higher Cr contents. The

coarsening of M23C6 can be reduced by adding B to the alloy but also by reducing the Cr

content in the composition, which is the main forming element of this precipitate.

Finally it must be pointed out that the designed alloys have relative large Co-contents due

to the high amount of ferrite stabilising elements. This is a drawback for the production

of low cost creep steels.

9% Cr alloys were designed to contain fine dispersed precipitates which present low

coarsening rates and are mostly placed within the sub-grains and at the sub-grain

boundaries for pinning of dislocations and/or sub-grains boundaries.

The designed 9% Cr alloys in this work showed promising creep strengths at 650°C / 100

MPa, reaching up to 15,000 h creep rupture lives. The alloys contain basically 2% W and

1% Co with a balanced content of B and N. They are similar to the Japanese 9Cr-3W-

3Co-B-N currently investigated by Abe et al. [24] but technically cheaper due to the

reduced amounts of Co and W compared to the Japanese steel.

As for 12% Cr alloys, best creep strength was obtained for the combination of

precipitates MX, M23C6 and Laves phase.


109

Ti-MX: Hardening of the alloy by precipitation of fine dispersed Ti-based MX particles

was achieved. Precipitation of large Ti-MX formed in the melt must be avoided because

they reduced considerably the creep strength. The precipitation of these carbides was

limited to the austenisation and tempering treatment used. For future investigations it

would be recommended to increase the austenisation temperature and to adjust the

tempering temperature in order to obtain a fine dispersion of this stable precipitates. As

discussed above, this fact means a challenge for industrial applications, where the

austenisation temperature should not be above 1150°C.

C-content: The effect of the carbon content on the 9% Cr alloys was investigated. The C

content affected directly the volume fraction of M23C6 carbides. Among all designed

alloys, the composition 9Cr-0.05C-1Co-2W-0.5Mn-0.35Si-0.035Nb-0.15V-0.01B-

0.005N showed best creep results. This 9% Cr alloy with reduced carbon content (~

0.05%) showed better distribution of M23C6 and MX precipitates, as well as minimal

coarsening of Laves phase after creep. Reducing the C content reduce the volume fraction

of M23C6, so that more nucleation sites for a fine dispersion of Laves phase are present.

This situation may favourable to avoid the formation of large Laves phases which are

detrimental for the creep strength of the alloy.

Summarising, future works on alloy design and production of 9-12% Cr heat resistant

steels for operations at 650°C / 100 MPa should consider the combination of creep

strength and oxidation resistance by adjusting the Cr content and by addition of alloying

elements in order to produce microstructures with precipitation of MX, M23C6 and Laves

phase, eventually Z-phase, for pinning of the sub-grain structure. Also a fine dispersion

and reduced coarsening of Laves phase may be intended.

References

110

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[5] C. Scheu, F. Kauffmann, G. Zies, K. Maile, S. Straub, and K. H. Mayer, “Requirements

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[8] P. D. Clarke, P. F. Morris, N. Cardinal, M. J. Worral, “Factors influencing the creep

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Curriculum Vitae Personal information

Last Name

First Name

Nationality

Date and place of birth

Education

Since September 2007

March 2000 to August 2007

Professional experience

Since September 2007

March 2007 to August 2007

January 2006 to March 2006

Rojas

David

Chilean

30th December 1980, Constitución, Chile.

Max-Planck Institute for Iron research, Düsseldorf (Germany) Department of Material Diagnostics and Steel Technology, PhD studies on “9-12% Cr heat resistant steels: alloy design, TEM characterisation of microstructure evolution and creep response at 650°C”. Engineering Studies at “Universidad de Concepción”, Concepción, Chile. Degree: Materials Engineering (Dipl.-Ing) Final Degree Project: “Study of the interaction slag/chloride salts in the production of secondary aluminium”.

PhD student at Max Planck Institut für Eisenforschung GmbH. (Germany). Teaching assistant of phase transformations Universidad de Concepción (Chile). Foundry simulation area. Foundry and machine shop Neptune (Iquique, Chile.)

9-12% Cr heat resistant steels : alloy design, TEM characterisation ...

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