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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.
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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
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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
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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
Metallurgy of 9-12% Cr steels
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:
Metallurgy of 9-12% Cr steels
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
Metallurgy of 9-12% Cr steels
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].
Metallurgy of 9-12% Cr steels
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.
Metallurgy of 9-12% Cr 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
Metallurgy of 9-12% Cr steels
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].
Metallurgy of 9-12% Cr steels
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.
Metallurgy of 9-12% Cr steels
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
Metallurgy of 9-12% Cr steels
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.
Metallurgy of 9-12% Cr steels
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].
Metallurgy of 9-12% Cr steels
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
Metallurgy of 9-12% Cr steels
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
Metallurgy of 9-12% Cr steels
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].
Metallurgy of 9-12% Cr steels
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
Metallurgy of 9-12% Cr steels
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)
Metallurgy of 9-12% Cr steels
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.
Results of 12% Cr Steels
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
Results of 12% Cr Steels
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].
Results of 12% Cr Steels
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).
Results of 12% Cr Steels
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].
Results of 12% Cr Steels
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.
Results of 12% Cr Steels
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
Results of 12% Cr Steels
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
Results of 12% Cr Steels
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.
Results of 12% Cr Steels
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).
Results of 12% Cr Steels
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.
Results of 12% Cr Steels
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.
Results of 12% Cr Steels
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.
Results of 12% Cr Steels
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.
Results of 12% Cr Steels
53
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.
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).
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.
Results of 12% Cr Steels
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.
Results of 12% Cr Steels
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.
Results of 12% Cr Steels
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.
Results of 12% Cr Steels
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.
Results of 12% Cr Steels
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
Results of 12% Cr Steels
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).
Results of 12% Cr Steels
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.
Results of 12% Cr Steels
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.
Results of 12% Cr Steels
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)
Dislocation density ρ (1013 m-2)
Scatter in ρ (1013 m-2) HV10
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
Results of 12% Cr Steels
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.
Results of 12% Cr Steels
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)
Dislocation density ρ (1013 m-2)
Scatter in ρ (1013 m-2) HV10
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.
Results of 12% Cr Steels
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
Discussion of 12% Cr Steels
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
Discussion of 12% Cr Steels
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].
Discussion of 12% Cr Steels
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].
Discussion of 12% Cr Steels
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].
Discussion of 12% Cr Steels
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.
Discussion of 12% Cr Steels
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
Discussion of 12% Cr Steels
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.
Discussion of 12% Cr Steels
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
Discussion of 12% Cr Steels
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.
Discussion of 12% Cr Steels
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.
Discussion of 12% Cr Steels
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.
Results of 9% Cr Steels
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.
Results of 9% Cr Steels
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.
Results of 9% Cr Steels
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.
Results of 9% Cr Steels
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
Results of 9% Cr Steels
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.
Results of 9% Cr Steels
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
Results of 9% Cr Steels
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.
Results of 9% Cr Steels
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
Results of 9% Cr Steels
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.
Results of 9% Cr Steels
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).
Results of 9% Cr Steels
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).
Results of 9% Cr Steels
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
Results of 9% Cr Steels
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).
Results of 9% Cr Steels
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.
Results of 9% Cr Steels
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
Results of 9% Cr Steels
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.).
Results of 9% Cr Steels
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.
Results of 9% Cr Steels
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
Results of 9% Cr Steels
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)
Results of 9% Cr Steels
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.
Results of 9% Cr Steels
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).
Discussion of 9% Cr steels
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).
Discussion of 9% Cr steels
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
Discussion of 9% Cr steels
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
Discussion of 9% Cr steels
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
Discussion of 9% Cr steels
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.
Discussion of 9% Cr steels
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.
Discussion of 9% Cr steels
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.
Final conclusion and perspectives
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.
Final conclusion and perspectives
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.
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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.)