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Microstructural study of creep rupture in a 12% chromium ferritic steel
Acra merail. Vol. 37, No. I, pp. 49-60, 1989 Printed in Great
Britain.
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All rights reserved
c
0001-6160189 $3.00 + 0.00 1989 Pergamon Press plc
MICROSTRUCTURAL STUDY OF CREEP RUPTURE IN A 12% CHROMIUM FERRITIC STEEL G. EGGELER,
J. C. EARTHMAN, N. NILSVANG and B. ILSCHNER Laboratoire de Metallurgic Mecanique, Ecole Polytechnique Fed&ale de Lausanne, Chemin de Bellerive 34, CH-1007 Lausanne, Switzerland (Received 18 December
1987; in revised form20 April 1988)
Abstract-Creep cavitation and rupture of a 12% Cr-M-V steel has been studied using advanced analytical techniques. By electronically superposing etched and unetched micrographs of cavitated regions, it has been possible to correlate cavity size and form to location within the complex microstructure of this material. The results indicate that cavitation occurs not only on former austenite grain boundaries (FAGBs) but also within the former austenite grains. The analysis also indicates, however, that the cavities lying on some of those FAGBs which are nearly perpendicular to the applied stress are larger and closer spaced than those within the grains and on other FAGB facets. Calculations of cavity size using a constrained cavity growth model are shown to agree well with the cavity size measurements for three different strains. The success of the constrained cavity growth description implies that cavitating facets in 12% Cr-Mo-V steel quickly shed load to other regions of the microstructure and thus act like microcracks long before cavity coalescence beomes evident. The possible role of sulphur in the nucleation of cavities has also been investigated. From Auger mapping results, it is concluded that larger concentration of sulphur are typically present where cavity nucleation takes place. This suggests that sulphur facilitates cavity nucleation in this material. R&urn&-La cavitation et la rupture d’un acier a 12% Cr lors du fluage ont et6 Ctudies par des methodes analytiques modernes. La superposition inforrnatisie de micrographics obtenues par microscopic optique et tlectroniques (MEB) de regions contenant des cavites a permis de corrtler leurs caracteristiques geombtriques en fonction de leur position au sein de la microstructure complexe de ce mattriau. Les observations indiquent que la cavitation se produit non seulement sur les joints de anciens grains austenitiques mais igalement dans I’inttrieur des anciens grains austinitiques. Neanmoins les cavites portees par des anciens joints de grains austlnitiques perpendiculaires a I’axe de traction, tendent a etre plus grandes que les autres cavitb. I1 a ete montre, que la croissance des cavitts est limit&e par le fluage de la matrice (“constrained cavity growth”). Les resultats numtriques obtenus a I’aide de ce modtle Concorde bien avec les dimensions mesurees des cavites. Le succ& de cette description permet d’inferer une rapide chute de contrainte sur les faces des anciens grains austinitiques portant des cavitis. Cela rend leur comportement analogue a celui de microfissures bien avant que ne se produise une coalescence gentrahste des cavites. Une analyse par sonde Auger a revBIt une plus forte concentration en soufre dans les cavites. Cela suggere que le soufre facilite la germination des cavites dans ce mattriau. und Kriechbruch eines ferritischen Stahles mit 12% Cr (deutsche Bezeichnung: X22 CrMoV 12 1) werden mit fortgeschrittenen gefiigeanalytischen Methoden untersucht. Durch rechnergestiitztes Uberlagern ungeatzter und geatzter Schliffbilder von porengeschadigten Bereichen war es moglich. Porengrosse und -form in Abhangigkeit von der jeweiligen Position der Pore im komplexen Mikrogefiige des Werkstoffs zu beschreiben. Porenbildung wird nicht nur auf ehemahgen Austenitkorngrenzen, sondern such im Inneren der ehemaligen Austenitkiirner beobachtet. Wie die quantitative Analyse zeigt, sind Poren auf solchen ehemaligen Austenitkorngrenzabschnitten, die senkrecht zur Spannungsachse liegen, grosser und hegen enger aufgereiht als andere Kriechporen. Gemessene Porengrossen stimmen gut mit errechneten Porengrossen iiberein, wenn man annimmt, dass Porenwachstum nicht unabhlngig vom Kriechen der umgebenden Matrix erfolgen kann (constrained cavity growth). Durch Porenbildung auf ehemaligen Austenitkorngrenzen kommt es zu einer Spannungsumverteilung im Mikrogefiige. Porengeschadigte frtihere Austenitkorngrenzen wirken bereits als Mikrorisse, bevor die Poren andgiiltig zusammengewachsen sind. Aus Untersuchungen mit der Augersonde kann gefolgert werden, dass dem Schwefel eine entscheidende Rolle bei der Porenkeimbildung zukommt. Zusammenfassung-Porenbildung
1. INTRODUCTION
Nine to twelve percent chromium ferritic steels are used for components in power generating and chemical industries, operating under creep conditions [ 1,2]. Currently, newly developed chromium steels with improved creep strength are in the focus of interest [3], and there are efforts in improving meth-
ods for predicting residual lifetime for established versions [4, 51. Often alloy development and design as well as residual lifetime estimation are based on creep rupture data alone. In order to progress in each of these areas, a better understanding of the role of microstructure during creep deformation and rupture is important. 49
50
EGGELER
et al.:
MICROSTRUCTURAL
STUDY OF CREEP RUPTURE
Advanced micromechanical creep rupture modelling for complex structural materials also requires and i.e. a description input, metallurgical quantification of creep damage development. This study concentrates on the microstructural features of creep rupture in a 12% chromium ferritic steel. Since at high temperatures crystalline solids can fracture by different mechanisms at different stress levels [26], high temperature rupture has been studied for stresses of 175 and 80 MPa at 923 K. Special emphasis is placed on the development of creep damage at the low stress level. There, cavity formation is studied focussing on both cavity nucleation and growth. Cavity size distributions were measured by means of a fully automatic image analyzing technique, which has been specially developed for high chromium ferritic steels. This technique takes into account the position of cavities relative to other microstructural elements like former austenite grain boundaries. Inhomogeneous creep cavitation is shown to be an essential feature of the damage accumulation process. This study is aimed at contributing to the general discussion of cavity nucleation, cavity growth and creep damage accumulation, based on microstructural observations in a heat resistant steel after creep testing. 2. EXPERIMENTS The composition of the 12% Cr-MO-V steel used in this work is shown in Table 1. The heat treatment of the as-received material consists of austenitizing at 1323 K for one hour followed by air cooling and tempering at 1023 K for four hours which is also followed by air cooling. Recent quantitative metallographic studies using transmission electron microscopy have revealed that the tempered martensite structure which results from this heat treatment basically consists of an elongated dislocation network that develops during tempering within the former austenite grains [6-91. The subgrain boundaries of this dislocation network stabilized by carbides, and high angle ferrite boundaries represent internal surfaces, which exist within the former austenite grains (FAGBs). Figure l(a) shows a high magnification transmission electron micrograph of the as received structure. A subgrain boundary intersects a FAGB which contains closely spaced Me,,C, carbides. Subgrains and carbides in the tempered martensite structure are too small to be resolved in the optical micrograph of Fig. l(b). Uniaxial constant load creep tests were performed at 923 K. Details of the creep testing are described elsewhere (9,101. The creep rupture data of this Table I. Chemical
composition
of the type X22Cr-MO-V (wt%)
12 I steel
C
Si
Mn
P
s
Cr
Ni
MO
V
0.24
0.49
0.68
0.017
0.005
12.04
0.72
1.08
0.29
Fig. 1. (a) TEM micrograph of the as-received material’s structure. A subgrain boundary intersects a former austenite grain boundary which contains closely spaced Me,,C, carbides. (a) Optical micrograph of the as-received 12% Cr-MO-V steel. are plotted in Fig. 2(a) as time to rupture vs stress on a log-log scale. Some of the corresponding continuously measured creep curves are presented in Fig. 2(b) where the logarithm of the strain rate is plotted as a function of strain. As can be seen in Fig. 2(c), the creep ductility which is represented by the elongation to rupture decreases with decreasing stress. “Creep brittleness” at low stresses has also been found in long time tests at typical service temperatures Ill], and was also observed for a ferritic steel with a lower chromium content [12]. Creep rupture at 923 K was studied for two constant load conditions (A and B in Fig. 2): condition A where the initial stress was 175 MPa and condition B where the initial stress was 80 MPa. Three of the creep tests performed under an initial stress of 80 MPa were terminated after 1. 5 and 12% respectively, with the objective to study the development of cavitation during creep; three tests were terminated by the rupture of the specimen. The rupture data for these three tests are given in Table 2. Structural investigations were performed by means of optical microscopy (OM), scanning electron microscopy (SEM) and Auger electron spectroscopy investigation
EGGELER
36
3.6 -2
et rrl.:
MICROSTRUCTURAL
360 3600 rupture time/loos
36ooO
STUDY
OF CREEP
RUPTURE
51
360000
-3
1
Fig. 3. SEM micrograph of the rupture surface of a specimen tested at 923 K and 175 MPa (labelled A in Fig. 2). strain
electron 1 pm.
o.50)
beam used for the AES analysis was approx.
3. RESULTS
3.1. Rupture surfaces at the high and low stress leael
o-l, 0
I,
60
60
I
100
120
I
I1
140
160
8’ 180
200
stress/MPa Fig. 2. Creep behavior of a 12% Cr-MO-V ferritic steel in constant load creep tests at 923 K: (a) time to rupture vs initial stress on a log-log scale; (b) continuously measured creep curves plotted as logarithm of strain rate vs true strain; (c) elongation to rupture vs initial stress.
(AES). Distributions of cavity features for different kinds of cavities were obtained using a quantitative metallographic method resolving cavities as small as 0.1 pm, which has been described elsewhere [13]. The chemical compositions of inclusions which were present in the as-received material were analyzed in the SEM with an energy dispersive X-ray spectrometer (EDAX). The rupture surfaces of the specimens were studied in the SEM. A compositional study on a cross-section of crept material was performed using a Perkin-Elmer 550AES. Surface contamination due to specimen preparation was removed by 2 keV Ar+ ion sputtering before determining the composition with AES. Sputtering was kept at a minimum (a few atomic monolayers) to avoid the removal of atoms which may have segregated at internal surfaces during heat treatment and creep. The diameter of the primary Table 2. Creep rupture
data from three creep tests at 923 K for an initial stress of 80 MPa Test
Rupture time (s) Strain to ruoture
(%)
A SEM micrograph of the rupture surface of a specimen tested at 923 K under the high stress condition (175 MPa, labelled A in Fig. 2) is shown in Fig. 3. For this stress level, the strain to rupture is 40% and specimens fail in a ductile manner characterized by large, sharply defined dimples containing second phase particles (Fig. 3). EDAX measurements indicate that these particles, which have diameters ranging from about 5 to 40pm, are oxides and manganese sulfides. These inclusions are also observed in the as-received material and, therefore. result from the processing of the steel. A rupture surface of an experiment under an initial stress of 80 MPa (labelled B in Fig. 2) is shown in Fig. 4. This specimen failed at a much smaller strain, I3%, and its rupture surface features are completely different from those of the specimen tested under the higher stress. No inclusions are observed on the
I
7.34 x 10s 12
Test 2
Test 3
8.82 x IO’ 8
1.36 x lob 14
Fig. 4. SEM micrograph of the rupture surface of a specimen tested at 923 K and 80 MPa (labelled B in Fig. 2).
52
EGGELER
et al.:
MICROSTRUCTURAL
STUDY OF CREEP RUPTURE
rupture surface and it is not possible to discern former austenite grain boundary facets. The size and shape of the granular features in Fig. 4 correspond to that of larger subgrains which were observed using TEM [9]. EDAX analysis of these small grains show the composition of the matrix material only. Cross-sections perpendicular to this rupture surface are shown in Fig. 5(a) and (b) (the applied stress direction is indicated). The white zone between the specimen and the mounting material is a nickel coating which was deposited electrochemically. Figure 5(a) shows that large cracks have formed perpendicular to the stress axis. A micrograph of the region near the rupture surface is shown in Fig. 5(b). A FAGB segment is indicated by an arrow. The
Fig. 6. Dark field optical micrographs. (a) As received material. (b) Material after 1% creep strain at 923 K and an initial stress of 80 MPa.
rupture surface only partially follows the FAGBs. Note that there is a FAGB segment that lies parallel to the rupture surface and which contains no cavities. 3.2, Development of damage during creep at the low stress level
Fig. 5. Optical micrographs of a metallographic cross section perpendicular to the rupture surface of a specimen tested at 923 K and 80 MPa (labelled B in Fig. 2). (a) Low magnification: cracks perpendicular to the stress axis. (b) High magnification: crack path and uncavitated FAGBs perpendicular to the stress axis close to the rupture surface.
Optical micrographs of the as-received material and material from a specimen which was crept to a final strain of 1% at 923 K and 80 MPa are shown in Fig. 6(a) and (b) respectively. The micrographs were taken using dark-field illumination [33]. The bright spots on the micrograph of the as-received material correspond to the large oxide and sulfide inclusions. After 1% creep strain, many smaller spots appear. These spots correspond to creep cavities that formed in the early stages of creep deformation. A small cavity after 5% creep strain is shown in the SEM micrograph in Fig. 7. A particle lying within the cavity is probably a Me,,C, carbide since it has a typic1 carbide size and shape as observed using TEM [Fig. l(a)]. Creep tests at 923 K and 80 MPa were terminated after 1, 5 and 12% strain to study the development of creep cavitation. A cross-sectional area of 0.2 pm* was chosen as the smallest cavity size for statistical evaluation for the 5 and 12% strain states. A smaller minimum size limit was chosen for the 1% strain state
EGGELER et al.:
MICROSTRUCTURAL
53
STUDY OF CREEP RUPTURE
Fig. 7. SEM micrograph of a cavity after 5% creep at 923 K with an initial stress of 80 MPa. -1
to obtain a sufficient sampling of the cavity size distribution. Cavity information was evaluated from sampling fields bounded by microhardness indentations. A SEM micrograph of the sampling fields for the specimen that crept to a strain of 12% is shown in Fig. 8. Note that the cavitation is heterogeneous: there are fields with high and low cavity densities. Two microcracks in one of the more cavitated fields. These were not included in the cavity measurements. Logarithmic cavity size distributions are shown in Fig. 9 for different creep strains at 923 K and 80 MPa. The mean value of logarithmic cavity size values in Fig. 9, increases with strain. A plot of the total cavity density (represented by the numbers of cavities per evaluated cross-section area) vs time is shown in Fig. IO(a) and a corresponding plot of total cavity density vs strain is illustrated in Fig. 10(b). The correlation between strain and cavity density in Fig. IO(b) appears to be linear while the correlation between time and cavity density is non-linear, Fig. IO(a). Cavities are not only found on the former austenite grain boundaries but also within the former austenite
-0.5 log
0.5
0 (cavity
area/pm?)
Fig. 9. Cumulative cavity size distributions for the total cavity population for different creep strains at 923 K and 80 MPa.
grains, as shown in Fig. 11 for a specimen crept to 5% strain. For this specimen, cumulative cavity size distributions for three different microstructural locations are shown in Fig. 12. Cavities on FAGBs and triple points tend to be larger than those in the interior of the former austenite grains. The cavity density, given by the number of cavities per evaluated
(a)
200
100 t
in
h
/, ---
lb)
5 E
Fig. 8. SEM micrograph of a 200 x 3OOpm’ area which is subdivided in six sampling fields by microhardness indentations.
in
ai0
Fig. 10. Development of the total cavity density during creep. (a) Cavity density as a function of time. (b) Cavity density as a function of strain.
EGGELER
54
er al.:
MICROSTRUCTURAL
Fig. 11. Optical micrograph of cavitated 12% Cr-MoV steel. Cavities are not only found on former austenite grain boundaries (FAGBs), but also in the interior of the former austenite grains.
cross-section area, is high for the cavities in the interior of the former austenite grains (4.3 x I03mme2) and on FAGBs (1.9 x 103mmm2) and low for the cavities on triple points (5.7 x 10-2mm-2). In the following, special emphasis is given to the cavities on the FAGBs. It is obvious that these cavities play a special role in the rupture process, not only due to their density and size, but also because they occur in planar arrays along FAGBs (Fig. 13). The effect of the angle fi between the grain boundary segment and the stress axis is illustrated in Fig. 14(a). Cavities are observed on grain boundaries of all orientations with respect to the stress axis. The mean cavity size is larger, however, for the cavities on FAGBs with angles between 60 and 90” with respect to the stress axis. Considering the line-up of cavities of FAGBs a parameter @ was evaluated from the test
STUDY OF CREEP RUPTURE
Fig. 13. Optical micrographs of cavitated 12% Cr-Mo-V steel. Cavities line up along FAGB segments perpendicular to the stress axis. region for the same three angle classes, which is given by the cavity number in the sampling field normalized by the total length of FAGB segments in the corresponding angle class and the number of FAGB segments. (This parameter gives a higher value for 99
f 2
16
2 z 14
0.7
I 0.3
I 0.1
0.6
I 0.9
log (cavity area/pm’)
(b)
boundaries
0' Creep:
0.3
T = 65OT, d=60N/mmz, 0.1
0.5
log (cavity area/
0.9
60'
20' P
~5% 1.1
Pm*)
Fig. 12. Cumulative cavity size distributions for cavities on triplepoints, on FAGB segments and in the interior of the former austenite grains after 5% strain. The scale on the vertical axis corresponds to an inverse Gause function.
Fig. 14. (a) Cumulative FAGB cavity size distributions for three classes of angle between the corresponding grain boundary segment and the stress axis for 5% strain. The scale on the vertical axis corresponds to an inverse Gauss function. (b) Parameter @ quantifying the arrangement of cavities on FAGB segments for three classes of angle between the correspnding grain boundary segments and the ,_^1 stress axes (3% stram).
EGGELER
Fig. 15. Optical form on FAGB
et d.:
micrograph of FAGB segments perpendicular (5% strain).
MICROSTRUCTURAL
microcracks that to the stress axis
three cavities on one FAGB segment of 30 pm length than for three cavities on three FAGB segments of 10 pm length.) A plot of @ for the three angle classes in Fig. 14(b) indicates that Q, also increases with increasing angle between the grain boundaries and the stress axis. Closely spaced FAGB cavities such as those on the optical micrograph in Fig. 13 eventually coalesce and form microcracks. Microcracks like those in Fig. 15 are rare: they are heterogeneously distributed throughout the test regions. In this study no effort has been made to study microcrack densities, because the number of microcracks observed after 5% deformation was insufficient to be treated statistically.
Fig. 16. Sulfur
distribution mapping
STUDY
OF CREEP
RUPTURE
55
Cross-sections of cavitated material crept to 5% strain were analyzed in the Auger spectrometer using a 1 firn beam diameter. Auger spectra inside cavities show pronounced sulfur peaks, which are not observed in energy spectra taken from uncavitated material. Sulphur detected within the cavity was probably present during the prior creep experiment. The sulfur distribution of the cavitated material as obtained by Auger mapping is shown in Fig. 16. The mean spacing of the bright regions, indicating the presence of sulphur, corresponds to the mean cavity spacing. This spacing (approx. 20 pm, Fig. 16) is much larger than the carbide spacing (approx. 0.5 pm [9]) and significantly smaller than the mean distance between the large sulfide inclusion [approx. 200 pm, Fig. 6(a)]. The size of the bright regions is greatly influenced by the geometry of the cavities and does not represent the actual cavity size. 4. DISCUSSION
OF EXPERIMENTAL
RESULTS
4. I. Ductile rupture at high stress levels High temperature rupture in a 12% Cr-MO-V steel has been studied for two stress levels at 923 K. At the high stress level, the specimens rupture in a ductile manner where voids form and grow plastically at large inclusions, which act as hard particles in a plastically deforming matrix. In a recent study of ductile rupture at room temperature of a ferritic steel with 9% chromium [14], voids form a MezlC, particles. Void formation was not observed until strains that reach and exceed the rupture strain of the
on cavitated 12% Cr-Mo-V steel after 5% creep deformation with an energy window around the sulfur KLL energy).
(AES sulfur
56
EGGELER et al.: MICROSTRUCTURAL STUDY OF CREEP RUPTURE
12% Cr-MO-V steel investigated in this work. For the high temperature experiments conducted in this investigation, large oxide and sulfide inclusions (5540 pm, Fig. 3) are more effective in the initiation of voids which grow by plastic deformation than small Me,,C, particles [0.1-0.3flm, Fig. l(a)]. The nature of plastic void formation processes at hard particles does not seem to change at high temperatures for high stresses. It is concluded that with increasing temperatures smaller particles become less effective in plastic void formation and do not play a role in the ductile high temperature rupture of 12% Cr-MO-V steel. 4.2. Cavity nucleation At the low stress level, diffusive creep cavitation appears to be the dominant damage mechanism which results in much lower strains to rupture. Cavities are observed as early as 1% strain [Fig. 6(b)] during primary creep where the strain rate is still decreasing towards the minimum creep rate [Fig. 2(b)]. Middleton [15] observed cavitation in iron based alloys as early as 0.05% strain. It is probable, therefore, that an incubation strain for nucleation is either small or not required for the cavities that form first. Nucleation of cavities occurs continuously during creep and the cavity density is proportional to strain [Fig. 10(b)]; this is also true for other engineering materials [16, 171. If nucleation theory by vacancy condensation is basically correct, it must be concluded that continuous cavity nucleation is a consequence of some stress concentration process which occurs throughout creep life and which triggers cavity nucleation. Argon et al. [27] propose that the triggering event for cavity nucleation might come from the spasmodic nature of grain boundary sliding. It appears unlikely that grain boundary sliding is strongly involved in cavity nucleation for the 12% Cr-M-V steel investigated in this work, since the density of Me,,C, carbides on FAGBs is very high [Fig. l(a)]. The absence of extensive grain boundary sliding is also indicated by the fact that triple point cavities were never found to be wedgecrack like, as one would expect them to be if they were merely a result of grain boundary sliding. Moreover this model cannot explain how cavities are nucleated on transverse FAGBs which do not slide. A more likely mechanism without grain boundary sliding has been proposed by Dyson [16]. It involves the intersection of slip bands with grain boundaries. This mechanism was observed by Nieh and Nix [19] for copper and by Shiozawa and Weertman [29] for a nickle base superalloy, where they found an almost perfect one-to-one correlation between the cavity spacing and the spacing between coarse slip bands. During creep of 12% Cr-Mo-V steel, slip bands do not form [9]. Thus, it is unlikely that cavity nucleation in 12% Cr-Mo-V steel occurs according to the model proposed by Dyson [16].
The AES study conducted in this work suggest that sulphur facilitates creep cavitation in the 12% Cr-Mo-V steel. This interpretation is supported by recent findings of George et al. [ 181. They reported, that for sulfur concentrations above the solubility limit (> 30 ppm in their study), sulphur can segregate between a particle and the surrounding matrix. Once there, it weakens the interface and facilitates cavity nucleation. SEM micrographs like that in Fig. 7 in addition to the Auger results suggest that sulphur segregates at interfaces between the Me,,C, carbides and matrix. This segregation process is certainly dependent on the heat treatment of the as-received material and might also continue during the creep process itself. Assuming that for a given particle size there is a critical concentration of sulphur at a particle/matrix interface which has to be achieved for cavity nucleation, continuous cavity nucleation might also be triggered by continuous segregation. It is also likely that sulphur segregation is only one prerequisite for cavity nucleation, while the nucleation process only starts if other conditions are met as well. Thus the nucleation process also depends on the crystallographic details of grain boundaries [22,23], which can explain why not all boundaries contain cavities. Moreover, if cavitation occurs nonuniformly in the solid, local multiaxial stress states can develop even in an uniaxial creep experiment. This local multiaxial stress states also can influence cavity nucleation and growth during creep [20,21]. 4.3. Location of cavities in the microstructure
In this study cavities are not only found on FAGBs, but also in the interior of the former austenite grains. This is due to the existence of other internal surfaces which form during tempering of the martensitic structure. These are high angle ferrite boundaries and subgrain boundaries [Fig. l(a)], which can also act as segregation sites for sulphur and facilitate diffusion. Thus, cavities in the interior of the former austenite grains do not have to grow by power law creep alone, which might be expected for cavities growing within a grain [24], when no such internal surfaces are present. Since these boundaries intersect former austenite grain boundaries, they also provide transverse diffusion paths for cavity growth on longitudinal FAGBs, where a relative high number of cavities with appreciable size is observed. Although the applied stress would not drive diffusion along a longitudinal FAGB, the growth of such cavities could be facilitated by “atom plating” on intersecting ferrite and subgrain boundaries which are transverse to the stress axis. 4.4. Cavity growth and accumulation of creep damage Although the mean cavity size, evaluated from large specimen regions, increases linearly with creep strain, one should keep in mind that cavitation occurs
jnhomogeReously in the structure. There are sampling fields in Fig. 8 with high and hi cavity densities. Onfy one of the six sam$ing fields wontons two microcracks. Cavitation is a&o i~~omogen~us in the sense that onfy smm of the FAGi% perpendicufar to the stress axis ~~~t~~n cavities. In Rg. 5fb), for example, a FAGB segment near the fracture surface is shown that did not cavitate. It folfows from tbc quantitative ~va~~~tjon preened in Fig. $2, that cavities on FAGBs have a more ~rn~o~ta~t role in the rupture process. They are significantly larger than the caviiies within the former anstenite grtins and are arranged on FAGB facets. The @-parameter @p.mntifying cavity line up along FAG& in this study) is highest and cavity growth is fas@st on FAGB seg~ men@ which are nearly ~er~nd~~~~a~ to the a~~~~~d stress (Fig. 14).
ary facets which are transverse to the apptied stress.
Once a criticat number aE microcracks have coaksced, the situating is no longer one of hon~o~e* neous creep damage but one of creep crack grow&, and the stram rate ia the tertiary crefzp range is strmgiy influenced by the creep damage, The transition from secondary to tertiary creep strongly de* pen& on the statistical positioning 5i” the micros cracks, which are ~etero~eneons~y distribute fRg. 8). This explains why the time to rupture can deviate by a factor of two for the same ~x~e~~me~lta~conditions (see Tabie 2).
A mode1 of the constrained cavity growth process, deposed by Rice [30], is shown sc~~~tj~a~i~ in Fig, li”. A cavitating grain boundary facet ~~i~rneter d) is embedded in creeping body under an apphed stress, 6. The cavities a8 have a radius equal to c and D untform spacing, a. In an extension of the o~~j~~ treatment I3ff, the equation for the cavity growth rate is given as
fn a&r $0 ~~rn~r~ the afo~~~~t~oned wsu&s with theoreticd predictions, we will now study the applicability of the constrained eavity grawth theory, first ~~~~du~ed by Dysan [Zgj. As shown in Fig, 13, ca~~~~t~~~g facets in KY%@r-Me-V steel we t;wpicafIy isolated with respect to other cavitating gram boundaries. Similar&, the constrained cz~&y growth mode1 describes a rn~~~ost~Gtnr~ where isolated transverse grain boundary facets cavitate while the neighboring Kn equation (31, !P is the tip angie of the cavities bo~~~~~~s are und~m~~~. In this easi;r, the cavity d~~~~d by CQSJI= ~~~~~~where yb is the grain bound* growth rate wilt he limited by dislocation creep in the ary free. energy anld ySis the surface fry%energy. The surrounding matrix. sit&ring stress, a,, in equation (1) is equaf to ~~servat~o~~s of cavities early in the He of speci- @y, sin@)jZ_ Vames of the mate~a~ constants used to mens suggest that some of the iasger grain ~und~~~ determine the cavity growth rate using equation {t) cavities nucfeate upon or shortly after loading. As a are given in Table 3. Equation (I} describes thli first ~~~ro~irnat~on, thsr preferred cav~t~t~~~ on trans- growth of qu~i~e~~i~~b~urn cavities as opposed to XW-N grain boundary facets corresponds to the largest flatal crack-like cavities, ff the surface dit%sion is cavities observed in the 12% Cr-Mo-V steef investisu%cientIy SIOW,then the morphotogy of the cavities &a&d here. It fohews that the f~rrn~t~o~ of the first wil$ be Bat arzd crack-Eke, OW ob~er~~~io~~ of cmmicrooraeks could involve these largest cavities, whieb ities in 12% Cr-MeV st&, e.g. Fig. 7, indicais_?that nucleate shortly after loading and be os1grain boundcavities are not crack-like aad typicalXy grow in 8
EGGELER
58
et al.:
MICROSTRUCTURAL
Table 3. Physical constants n
B
(MPa)-66
used in the constrained
Power law stress exponent (650°C)
9.12 x 1O-2’
Power law creep coefficient (650°C)
(m’)
I.18 x 10-29”~21 Atomic
SD,
(m) SF’)
,,I x lO-‘W?
Pre-exponential for grain boundary diffusion
Qb
(J n-m-‘)
1.74 x 105”.*1
Activation energy for grain boundary diffusion
Ys
(J m-*)
2.1’2’
Surface free energy
Yb
(J mW2)
0.8512]
Grain
Plasticityand Creepof
quasi-equilibrium manner in this material. From microstructural measurements on cavitated facets such as the one shown in Fig. 13, a mean value of 20pm was determined for d and mean values of 1 ranging from 5 to 7 pm were determined for grain boundary facets nearly transverse to the applied stress. To a good approximation, a cavity may nucleate and grow if the cavity radius is greater than 2y,/a. This value for 12% Cr-Mo-V steel under 80 MPa, 2yJa = 0.07 pm, was taken to be the initial cavity radius, c,,, in the model. Cavity size is calculated in the model by numerically integrating equation (1). For a comparison with the measured cavity size data, the cavity radius was calculated for an applied stress of 80 MPa and creep times corresponding to 1, 5 and 12% strain. The model predictions are presented in Table 4 where they are compared with the largest 5% of the cavity population. These measured values are used as an indication of the upper limit of the observed cavity sizes. The model successfully predicts cavity sizes for 1 = 5 pm which lie within the measured 5% range. The results for 1 = 7 pm are also good approximations of the maximum observed cavity size demonstrating that the cavity spacing need not be extremely accurate to obtain reasonable predictions. It is evident from these findings that the contained cavity growth model is an accurate description of the grain boundary cavitation in 12% Cr-MeV steel. An unconstrained cavity growth model [32], in which cavity growth is only limited by the grain boundary diffusion, predicts cavity coalescence (2c = 1 = 7 pm) after only 1.3 x lo5 s for the same material and conditions. Since cavities are required to
I
5 I2
Time (s) 8.6 x 10’ 105 1.0 x 106
7.0 x
of upper limit cavity size measurements model calculations
0.6 to 0.9 I.8 1.3 to 1.9
1.2 to
boundary
free energy
Deformation-Mechanism Maps, The Metalsand Ceramics,pp. 62-63. Pergamon Press,
(1987).
Measured cavity radius (rm) [largest 5%]
volume
and M. F. Ashby,
Oxford (1982). ‘21H Riedel , Fracture at High Temperatures,
Strain (%)
cavity growth model
6.6
R
[“H. J. Edward
Table 4. Comparison
STUDY OF CREEP RUPTURE
Calculated cavity radius (rm) 1=5flm 1=7pm 0.7 1.5 1.7
0.8 I.8 2.1
with
Appendix
A. Springer,
Berlin
be present on all facets, cavity coalescence should correspond to the time to rupture for unconstrained cavity growth. The corresponding observed rupture time (approx. 1 x 106s), however, is about an order of magnitude greater than 1.3 x lo5 s. 4.6. Constrained formation
cavity
growth
and
microcrack
There is an important implication of the constrained cavity growth model which has to do with the nature of the stresses acting on a cavitating facet. The facet stress, cr,,in Fig. 17, rapidly decreases early in the course of constrained cavity growth. In this way, the cavitating facet behaves mechanically similar to a traction-free microcrack long before cavity coalescence occurs. Thus, the time to cavity coalescence on isolated facets is not a good criterion for predicting the time to rupture. This is further supported by the fact that, although there is good agreement between the measured and the calculated cavity size values, the calculated time to cavity coalescence (2.8 x lo6 s for 1 = 5 pm) is greater than the observed rupture time by about a factor of 3. Closely spaced microcracks, where cavity coalescence had occurred, were observed late in the life of the specimen. Although cavity coalescence is not predicted for the average grain facet size of 20 pm, it could occur by constrained cavity growth on abnormally large grain facets since the cavity growth rate increases with increasing facet size, d, in equation (1). It is likely that, between two closely spaced cavitating facets, transgranular deformation and damage is accelerated such that the two cavitating boundaries behave as one late facet. In this way the effective facet diameter could be much larger than the grain size. The formation of a large microcrack by this process is shown schematically in Fig. 18. Two closely spaced microcracks can be seen in one of the sampling fields in Fig. 8 which together comprise an effective facet size of about 70 pm. When the value of d in the constrained cavity growth model is 70pm, cavity coalescence (c = J.) occurs earlier than the corresponding observed rupture time. The presence of
EGGELER et al.: MICROSTRUCTURAL STUDY OF CREEP Cavitating Facet Interaction
\
o
I
\
I
Fig. 18. Sequential schematic drawings of a possible mechanism for the formation of microcracks by constrained cavity growth on interacting grain boundary facets.
RUPTURE
59
gation can occur preferentially and which will facilitate diffusion. A quantitative metallographic evaluation of the cavity population showed that cavities on transverse former austenite grain boundaries play the dominant role in the rupture process. The constrained cavity growth model reasonably agrees with the observed growth of closely spaced cavities on “isolated FAGBs”, however, overestimates the rupture time when cavity coalescence is used as a rupture criterion. Therefore, cavity coalescence should not be used as a criterion for creep rupture for a 12% Cr-Mo-V steel. The rupture time is determined by the formation and growth of microcracks from closely spaced cavitating grain boundary facets. Acknowledgements-Financial support of the Swiss Federal Government within the framework of the European cooperation COST 505 is gratefully acknowledged. Joint support from the Swiss Commission for the Promotion of Scientific Research and Sulzer AG, represented by Dr P. P. Schepp and Dr B. Walser, under contract KWF 1512 is appreciated. The authors wish to thank Dr H. Mathieu from the Laboratoire de Metallurgic Chimique for kindly making available the Auger spectrometer and for his help in performing the measurements and Professor W. D. Nix for fruitful discussions during his stay in Lausanne on sabbatical leave from Stanford University.
REFERENCES closely spaced microcracks is therefore consistent with constrained cavity growth theory since the effective facet size is sufficiently large. Although it is not typical for a cavitating facet to interact with another, it appears that the formation and growth of microcracks from the facets which do interact have an important role in determining the life of the specimen. 5. SUMMARY AND CONCLUSION High temperature rupture of a 12% Cr-Mo-V steel was studied for two stress levels at 923 K. At the high stress level (175 MPa) rupture occurs in a ductile manner. Voids form at oxide and sulfide inclusions which act as hard obstacles in a plastically deforming matrix. At the low stress level (80 MPa) rupture is controlled by the nucleation and diffusional growth of creep cavities, the first of which nucleate upon loading. Auger results suggest that segregation of sulfur at a carbide/matrix interface facilitates cavity nucleation which also is dependent on grain boundary crystallography and local state of stress. A characteristic feature of creep cavitation in high chromium ferritic steels is that cavities are not only found on former austenite grain boundaries (FAGBs) but also in the interior of former austenite grains. There are additional internal surfaces such as subgrain boundaries and high angle ferrite boundaries, which have formed during tempering, where segre-
1. J. Z. Briggs and T. D. Parker, The Super 12% Cr Steels. Climax Molvbdenum. New York (1965). 2. B. Walser, d. Brezina’and T. Greiger, krch. Eisenhtitt. 50, 249 (1979). 3. F. Masuyama and H. Haneda, The First Int. Conf on Improved Coal-Fired Power Plants (Proc. Conf), Palo
Alto, California, EPRI (1986). 4. B. Walser Heat and Mass Transfer in Metallurgical Svstems (Proc. Con/X. Dubrovnik. Yueoslavia. 1979.
Hemisphere, Washington, D.C. (1981). 5. R. Sandstrom, S. Karlson and S. Modin, High Temp. Tech. 3, 71 (1985). 6. E. A. Little, D. R. Harries, F. B. Pickering and S. R. Keown, Metals Tech. 4, 205 (1977). 7. I. M. Park, T. Fujita and K. Asakura, Trans. Iron Steel Inst. Japan 20, 99 (1980). 8. J. W. Schinkel, P. L. F. Radermakers, B. R. Drenth and C. P. Scheepens, J. Hear Treat. 3, 237 (1984). 9. G. Eggeler, N. Nilsvang and B. Ilschner, Sreel Res. 58,
97 (1987). 10. G. Eggeler, B. Ilschner, P. P. Schepp and R. Zohner, Mater. Technik. 14, 187 (1986).
I I. H. Weber, Festigkeits- und Bruchverhalten bei hiiheren Temaeraluren. Vol. 2. DV. l-69. Verlag Stahleisen. Dusseldorf, F’. R. G. (1980). 12. B. J. Cane, Mechanical Behauiour of Materials-III. (Proc. Conf.), Cambridge, England. Vol. 2. pp. 1733182, 1979 Pergamon Press, Oxford (1980). 13. N. Nilsvang and G. Eggeler, Prukf. Mefallogr. 24, 321 (1987). 14. B. A. Senior, F. W. Noble and B. L. Eyre, Acra metall. 34, 1321 (1986). 15. C. J. Middleton, Metal Sci. 15, 154 (1981). 16. B. F. Dyson, Scripfa melall. 17, 31 (1983). 17. H. Riedel, Fracrure at High Temperatures. Springer,
Berlin (1987). 18. E. P. George, P. L. Li and D. P. Pope, Creep and
EGGELER
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et al.: MICROSTRUCTURAL
Fracture of Engineering Materials and Structures (Proc. Conf.), Swansea, England. pp. 169-182. Inst. of Metals,
London (1987). 19. T. G. Nieh and W. D. Nix, Scripta metall. 14, 365 (1980). 20. D. R. Hayhurst and F. A. Leckie, Mechanical Behuviour of Materials-IV. (Suppl). Stockholm, Sweden, (1983).
Pergamon Press, Oxford (I 983). 21. B. J. Cane, Advances in Fracture Research (Fracture 81), 1981, Vol. 3, (Proc. Co&), Cannes, France,
pp. 128551293. Pergamon Press, Oxford (1982). 22. J. Don and S. Mahjumdar, Acta metall. 34, 961 (1986). 23. T. Watanabe, Creep and Fracture of Engineering Materials and Structures, (Proc. Conf.), Swansea, England, 1987. pp. 1555168. Inst. of Metals, London (1987). 24. A. C. F. Cocks and M. F. Ashby, Scripta metall. 16,109 (1982).
STUDY OF CREEP RUPTURE
25. D. J. Gooch, Metal1 Sci. 16, 79 (1982). 26. M. F. Ashby, Fracture 1977, Vol. 1, pp. 1-14 (1977). 21. A. S. Argon, I-W. Chen and C. W. Lau. Creep-Fatigue-Environment Interactions, pp. 4685.
Am. Inst. Min. Engrs, New York (1980). 28. B. F. Dyson, Metals Sci. 10,349 (1976). 29. K. Shiozawa and J. R. Wertmann, Acta metall. 31, 993 (1983). 30. J. R. Rice, Acta metall. 29, 675 (1981). 31. H. Riedel, Mechanical Behaviour ofMaterials-IV, Vol. 1, pp. 67-78. Pergamon Press, Oxford (1985). 32. I.-W. Chen and A. S. Argon, Acta metall. 29, 1759 (1981). 33. For example on page 76 of: Metals Handbook, Vol. 9, Metallography and Microstructures. Am. Sot. Metals,
Metals Park, Ohio (1985).
Britain.
Copyright
All rights reserved
c
0001-6160189 $3.00 + 0.00 1989 Pergamon Press plc
MICROSTRUCTURAL STUDY OF CREEP RUPTURE IN A 12% CHROMIUM FERRITIC STEEL G. EGGELER,
J. C. EARTHMAN, N. NILSVANG and B. ILSCHNER Laboratoire de Metallurgic Mecanique, Ecole Polytechnique Fed&ale de Lausanne, Chemin de Bellerive 34, CH-1007 Lausanne, Switzerland (Received 18 December
1987; in revised form20 April 1988)
Abstract-Creep cavitation and rupture of a 12% Cr-M-V steel has been studied using advanced analytical techniques. By electronically superposing etched and unetched micrographs of cavitated regions, it has been possible to correlate cavity size and form to location within the complex microstructure of this material. The results indicate that cavitation occurs not only on former austenite grain boundaries (FAGBs) but also within the former austenite grains. The analysis also indicates, however, that the cavities lying on some of those FAGBs which are nearly perpendicular to the applied stress are larger and closer spaced than those within the grains and on other FAGB facets. Calculations of cavity size using a constrained cavity growth model are shown to agree well with the cavity size measurements for three different strains. The success of the constrained cavity growth description implies that cavitating facets in 12% Cr-Mo-V steel quickly shed load to other regions of the microstructure and thus act like microcracks long before cavity coalescence beomes evident. The possible role of sulphur in the nucleation of cavities has also been investigated. From Auger mapping results, it is concluded that larger concentration of sulphur are typically present where cavity nucleation takes place. This suggests that sulphur facilitates cavity nucleation in this material. R&urn&-La cavitation et la rupture d’un acier a 12% Cr lors du fluage ont et6 Ctudies par des methodes analytiques modernes. La superposition inforrnatisie de micrographics obtenues par microscopic optique et tlectroniques (MEB) de regions contenant des cavites a permis de corrtler leurs caracteristiques geombtriques en fonction de leur position au sein de la microstructure complexe de ce mattriau. Les observations indiquent que la cavitation se produit non seulement sur les joints de anciens grains austenitiques mais igalement dans I’inttrieur des anciens grains austinitiques. Neanmoins les cavites portees par des anciens joints de grains austlnitiques perpendiculaires a I’axe de traction, tendent a etre plus grandes que les autres cavitb. I1 a ete montre, que la croissance des cavitts est limit&e par le fluage de la matrice (“constrained cavity growth”). Les resultats numtriques obtenus a I’aide de ce modtle Concorde bien avec les dimensions mesurees des cavites. Le succ& de cette description permet d’inferer une rapide chute de contrainte sur les faces des anciens grains austinitiques portant des cavitis. Cela rend leur comportement analogue a celui de microfissures bien avant que ne se produise une coalescence gentrahste des cavites. Une analyse par sonde Auger a revBIt une plus forte concentration en soufre dans les cavites. Cela suggere que le soufre facilite la germination des cavites dans ce mattriau. und Kriechbruch eines ferritischen Stahles mit 12% Cr (deutsche Bezeichnung: X22 CrMoV 12 1) werden mit fortgeschrittenen gefiigeanalytischen Methoden untersucht. Durch rechnergestiitztes Uberlagern ungeatzter und geatzter Schliffbilder von porengeschadigten Bereichen war es moglich. Porengrosse und -form in Abhangigkeit von der jeweiligen Position der Pore im komplexen Mikrogefiige des Werkstoffs zu beschreiben. Porenbildung wird nicht nur auf ehemahgen Austenitkorngrenzen, sondern such im Inneren der ehemaligen Austenitkiirner beobachtet. Wie die quantitative Analyse zeigt, sind Poren auf solchen ehemaligen Austenitkorngrenzabschnitten, die senkrecht zur Spannungsachse liegen, grosser und hegen enger aufgereiht als andere Kriechporen. Gemessene Porengrossen stimmen gut mit errechneten Porengrossen iiberein, wenn man annimmt, dass Porenwachstum nicht unabhlngig vom Kriechen der umgebenden Matrix erfolgen kann (constrained cavity growth). Durch Porenbildung auf ehemaligen Austenitkorngrenzen kommt es zu einer Spannungsumverteilung im Mikrogefiige. Porengeschadigte frtihere Austenitkorngrenzen wirken bereits als Mikrorisse, bevor die Poren andgiiltig zusammengewachsen sind. Aus Untersuchungen mit der Augersonde kann gefolgert werden, dass dem Schwefel eine entscheidende Rolle bei der Porenkeimbildung zukommt. Zusammenfassung-Porenbildung
1. INTRODUCTION
Nine to twelve percent chromium ferritic steels are used for components in power generating and chemical industries, operating under creep conditions [ 1,2]. Currently, newly developed chromium steels with improved creep strength are in the focus of interest [3], and there are efforts in improving meth-
ods for predicting residual lifetime for established versions [4, 51. Often alloy development and design as well as residual lifetime estimation are based on creep rupture data alone. In order to progress in each of these areas, a better understanding of the role of microstructure during creep deformation and rupture is important. 49
50
EGGELER
et al.:
MICROSTRUCTURAL
STUDY OF CREEP RUPTURE
Advanced micromechanical creep rupture modelling for complex structural materials also requires and i.e. a description input, metallurgical quantification of creep damage development. This study concentrates on the microstructural features of creep rupture in a 12% chromium ferritic steel. Since at high temperatures crystalline solids can fracture by different mechanisms at different stress levels [26], high temperature rupture has been studied for stresses of 175 and 80 MPa at 923 K. Special emphasis is placed on the development of creep damage at the low stress level. There, cavity formation is studied focussing on both cavity nucleation and growth. Cavity size distributions were measured by means of a fully automatic image analyzing technique, which has been specially developed for high chromium ferritic steels. This technique takes into account the position of cavities relative to other microstructural elements like former austenite grain boundaries. Inhomogeneous creep cavitation is shown to be an essential feature of the damage accumulation process. This study is aimed at contributing to the general discussion of cavity nucleation, cavity growth and creep damage accumulation, based on microstructural observations in a heat resistant steel after creep testing. 2. EXPERIMENTS The composition of the 12% Cr-MO-V steel used in this work is shown in Table 1. The heat treatment of the as-received material consists of austenitizing at 1323 K for one hour followed by air cooling and tempering at 1023 K for four hours which is also followed by air cooling. Recent quantitative metallographic studies using transmission electron microscopy have revealed that the tempered martensite structure which results from this heat treatment basically consists of an elongated dislocation network that develops during tempering within the former austenite grains [6-91. The subgrain boundaries of this dislocation network stabilized by carbides, and high angle ferrite boundaries represent internal surfaces, which exist within the former austenite grains (FAGBs). Figure l(a) shows a high magnification transmission electron micrograph of the as received structure. A subgrain boundary intersects a FAGB which contains closely spaced Me,,C, carbides. Subgrains and carbides in the tempered martensite structure are too small to be resolved in the optical micrograph of Fig. l(b). Uniaxial constant load creep tests were performed at 923 K. Details of the creep testing are described elsewhere (9,101. The creep rupture data of this Table I. Chemical
composition
of the type X22Cr-MO-V (wt%)
12 I steel
C
Si
Mn
P
s
Cr
Ni
MO
V
0.24
0.49
0.68
0.017
0.005
12.04
0.72
1.08
0.29
Fig. 1. (a) TEM micrograph of the as-received material’s structure. A subgrain boundary intersects a former austenite grain boundary which contains closely spaced Me,,C, carbides. (a) Optical micrograph of the as-received 12% Cr-MO-V steel. are plotted in Fig. 2(a) as time to rupture vs stress on a log-log scale. Some of the corresponding continuously measured creep curves are presented in Fig. 2(b) where the logarithm of the strain rate is plotted as a function of strain. As can be seen in Fig. 2(c), the creep ductility which is represented by the elongation to rupture decreases with decreasing stress. “Creep brittleness” at low stresses has also been found in long time tests at typical service temperatures Ill], and was also observed for a ferritic steel with a lower chromium content [12]. Creep rupture at 923 K was studied for two constant load conditions (A and B in Fig. 2): condition A where the initial stress was 175 MPa and condition B where the initial stress was 80 MPa. Three of the creep tests performed under an initial stress of 80 MPa were terminated after 1. 5 and 12% respectively, with the objective to study the development of cavitation during creep; three tests were terminated by the rupture of the specimen. The rupture data for these three tests are given in Table 2. Structural investigations were performed by means of optical microscopy (OM), scanning electron microscopy (SEM) and Auger electron spectroscopy investigation
EGGELER
36
3.6 -2
et rrl.:
MICROSTRUCTURAL
360 3600 rupture time/loos
36ooO
STUDY
OF CREEP
RUPTURE
51
360000
-3
1
Fig. 3. SEM micrograph of the rupture surface of a specimen tested at 923 K and 175 MPa (labelled A in Fig. 2). strain
electron 1 pm.
o.50)
beam used for the AES analysis was approx.
3. RESULTS
3.1. Rupture surfaces at the high and low stress leael
o-l, 0
I,
60
60
I
100
120
I
I1
140
160
8’ 180
200
stress/MPa Fig. 2. Creep behavior of a 12% Cr-MO-V ferritic steel in constant load creep tests at 923 K: (a) time to rupture vs initial stress on a log-log scale; (b) continuously measured creep curves plotted as logarithm of strain rate vs true strain; (c) elongation to rupture vs initial stress.
(AES). Distributions of cavity features for different kinds of cavities were obtained using a quantitative metallographic method resolving cavities as small as 0.1 pm, which has been described elsewhere [13]. The chemical compositions of inclusions which were present in the as-received material were analyzed in the SEM with an energy dispersive X-ray spectrometer (EDAX). The rupture surfaces of the specimens were studied in the SEM. A compositional study on a cross-section of crept material was performed using a Perkin-Elmer 550AES. Surface contamination due to specimen preparation was removed by 2 keV Ar+ ion sputtering before determining the composition with AES. Sputtering was kept at a minimum (a few atomic monolayers) to avoid the removal of atoms which may have segregated at internal surfaces during heat treatment and creep. The diameter of the primary Table 2. Creep rupture
data from three creep tests at 923 K for an initial stress of 80 MPa Test
Rupture time (s) Strain to ruoture
(%)
A SEM micrograph of the rupture surface of a specimen tested at 923 K under the high stress condition (175 MPa, labelled A in Fig. 2) is shown in Fig. 3. For this stress level, the strain to rupture is 40% and specimens fail in a ductile manner characterized by large, sharply defined dimples containing second phase particles (Fig. 3). EDAX measurements indicate that these particles, which have diameters ranging from about 5 to 40pm, are oxides and manganese sulfides. These inclusions are also observed in the as-received material and, therefore. result from the processing of the steel. A rupture surface of an experiment under an initial stress of 80 MPa (labelled B in Fig. 2) is shown in Fig. 4. This specimen failed at a much smaller strain, I3%, and its rupture surface features are completely different from those of the specimen tested under the higher stress. No inclusions are observed on the
I
7.34 x 10s 12
Test 2
Test 3
8.82 x IO’ 8
1.36 x lob 14
Fig. 4. SEM micrograph of the rupture surface of a specimen tested at 923 K and 80 MPa (labelled B in Fig. 2).
52
EGGELER
et al.:
MICROSTRUCTURAL
STUDY OF CREEP RUPTURE
rupture surface and it is not possible to discern former austenite grain boundary facets. The size and shape of the granular features in Fig. 4 correspond to that of larger subgrains which were observed using TEM [9]. EDAX analysis of these small grains show the composition of the matrix material only. Cross-sections perpendicular to this rupture surface are shown in Fig. 5(a) and (b) (the applied stress direction is indicated). The white zone between the specimen and the mounting material is a nickel coating which was deposited electrochemically. Figure 5(a) shows that large cracks have formed perpendicular to the stress axis. A micrograph of the region near the rupture surface is shown in Fig. 5(b). A FAGB segment is indicated by an arrow. The
Fig. 6. Dark field optical micrographs. (a) As received material. (b) Material after 1% creep strain at 923 K and an initial stress of 80 MPa.
rupture surface only partially follows the FAGBs. Note that there is a FAGB segment that lies parallel to the rupture surface and which contains no cavities. 3.2, Development of damage during creep at the low stress level
Fig. 5. Optical micrographs of a metallographic cross section perpendicular to the rupture surface of a specimen tested at 923 K and 80 MPa (labelled B in Fig. 2). (a) Low magnification: cracks perpendicular to the stress axis. (b) High magnification: crack path and uncavitated FAGBs perpendicular to the stress axis close to the rupture surface.
Optical micrographs of the as-received material and material from a specimen which was crept to a final strain of 1% at 923 K and 80 MPa are shown in Fig. 6(a) and (b) respectively. The micrographs were taken using dark-field illumination [33]. The bright spots on the micrograph of the as-received material correspond to the large oxide and sulfide inclusions. After 1% creep strain, many smaller spots appear. These spots correspond to creep cavities that formed in the early stages of creep deformation. A small cavity after 5% creep strain is shown in the SEM micrograph in Fig. 7. A particle lying within the cavity is probably a Me,,C, carbide since it has a typic1 carbide size and shape as observed using TEM [Fig. l(a)]. Creep tests at 923 K and 80 MPa were terminated after 1, 5 and 12% strain to study the development of creep cavitation. A cross-sectional area of 0.2 pm* was chosen as the smallest cavity size for statistical evaluation for the 5 and 12% strain states. A smaller minimum size limit was chosen for the 1% strain state
EGGELER et al.:
MICROSTRUCTURAL
53
STUDY OF CREEP RUPTURE
Fig. 7. SEM micrograph of a cavity after 5% creep at 923 K with an initial stress of 80 MPa. -1
to obtain a sufficient sampling of the cavity size distribution. Cavity information was evaluated from sampling fields bounded by microhardness indentations. A SEM micrograph of the sampling fields for the specimen that crept to a strain of 12% is shown in Fig. 8. Note that the cavitation is heterogeneous: there are fields with high and low cavity densities. Two microcracks in one of the more cavitated fields. These were not included in the cavity measurements. Logarithmic cavity size distributions are shown in Fig. 9 for different creep strains at 923 K and 80 MPa. The mean value of logarithmic cavity size values in Fig. 9, increases with strain. A plot of the total cavity density (represented by the numbers of cavities per evaluated cross-section area) vs time is shown in Fig. IO(a) and a corresponding plot of total cavity density vs strain is illustrated in Fig. 10(b). The correlation between strain and cavity density in Fig. IO(b) appears to be linear while the correlation between time and cavity density is non-linear, Fig. IO(a). Cavities are not only found on the former austenite grain boundaries but also within the former austenite
-0.5 log
0.5
0 (cavity
area/pm?)
Fig. 9. Cumulative cavity size distributions for the total cavity population for different creep strains at 923 K and 80 MPa.
grains, as shown in Fig. 11 for a specimen crept to 5% strain. For this specimen, cumulative cavity size distributions for three different microstructural locations are shown in Fig. 12. Cavities on FAGBs and triple points tend to be larger than those in the interior of the former austenite grains. The cavity density, given by the number of cavities per evaluated
(a)
200
100 t
in
h
/, ---
lb)
5 E
Fig. 8. SEM micrograph of a 200 x 3OOpm’ area which is subdivided in six sampling fields by microhardness indentations.
in
ai0
Fig. 10. Development of the total cavity density during creep. (a) Cavity density as a function of time. (b) Cavity density as a function of strain.
EGGELER
54
er al.:
MICROSTRUCTURAL
Fig. 11. Optical micrograph of cavitated 12% Cr-MoV steel. Cavities are not only found on former austenite grain boundaries (FAGBs), but also in the interior of the former austenite grains.
cross-section area, is high for the cavities in the interior of the former austenite grains (4.3 x I03mme2) and on FAGBs (1.9 x 103mmm2) and low for the cavities on triple points (5.7 x 10-2mm-2). In the following, special emphasis is given to the cavities on the FAGBs. It is obvious that these cavities play a special role in the rupture process, not only due to their density and size, but also because they occur in planar arrays along FAGBs (Fig. 13). The effect of the angle fi between the grain boundary segment and the stress axis is illustrated in Fig. 14(a). Cavities are observed on grain boundaries of all orientations with respect to the stress axis. The mean cavity size is larger, however, for the cavities on FAGBs with angles between 60 and 90” with respect to the stress axis. Considering the line-up of cavities of FAGBs a parameter @ was evaluated from the test
STUDY OF CREEP RUPTURE
Fig. 13. Optical micrographs of cavitated 12% Cr-Mo-V steel. Cavities line up along FAGB segments perpendicular to the stress axis. region for the same three angle classes, which is given by the cavity number in the sampling field normalized by the total length of FAGB segments in the corresponding angle class and the number of FAGB segments. (This parameter gives a higher value for 99
f 2
16
2 z 14
0.7
I 0.3
I 0.1
0.6
I 0.9
log (cavity area/pm’)
(b)
boundaries
0' Creep:
0.3
T = 65OT, d=60N/mmz, 0.1
0.5
log (cavity area/
0.9
60'
20' P
~5% 1.1
Pm*)
Fig. 12. Cumulative cavity size distributions for cavities on triplepoints, on FAGB segments and in the interior of the former austenite grains after 5% strain. The scale on the vertical axis corresponds to an inverse Gause function.
Fig. 14. (a) Cumulative FAGB cavity size distributions for three classes of angle between the corresponding grain boundary segment and the stress axis for 5% strain. The scale on the vertical axis corresponds to an inverse Gauss function. (b) Parameter @ quantifying the arrangement of cavities on FAGB segments for three classes of angle between the correspnding grain boundary segments and the ,_^1 stress axes (3% stram).
EGGELER
Fig. 15. Optical form on FAGB
et d.:
micrograph of FAGB segments perpendicular (5% strain).
MICROSTRUCTURAL
microcracks that to the stress axis
three cavities on one FAGB segment of 30 pm length than for three cavities on three FAGB segments of 10 pm length.) A plot of @ for the three angle classes in Fig. 14(b) indicates that Q, also increases with increasing angle between the grain boundaries and the stress axis. Closely spaced FAGB cavities such as those on the optical micrograph in Fig. 13 eventually coalesce and form microcracks. Microcracks like those in Fig. 15 are rare: they are heterogeneously distributed throughout the test regions. In this study no effort has been made to study microcrack densities, because the number of microcracks observed after 5% deformation was insufficient to be treated statistically.
Fig. 16. Sulfur
distribution mapping
STUDY
OF CREEP
RUPTURE
55
Cross-sections of cavitated material crept to 5% strain were analyzed in the Auger spectrometer using a 1 firn beam diameter. Auger spectra inside cavities show pronounced sulfur peaks, which are not observed in energy spectra taken from uncavitated material. Sulphur detected within the cavity was probably present during the prior creep experiment. The sulfur distribution of the cavitated material as obtained by Auger mapping is shown in Fig. 16. The mean spacing of the bright regions, indicating the presence of sulphur, corresponds to the mean cavity spacing. This spacing (approx. 20 pm, Fig. 16) is much larger than the carbide spacing (approx. 0.5 pm [9]) and significantly smaller than the mean distance between the large sulfide inclusion [approx. 200 pm, Fig. 6(a)]. The size of the bright regions is greatly influenced by the geometry of the cavities and does not represent the actual cavity size. 4. DISCUSSION
OF EXPERIMENTAL
RESULTS
4. I. Ductile rupture at high stress levels High temperature rupture in a 12% Cr-MO-V steel has been studied for two stress levels at 923 K. At the high stress level, the specimens rupture in a ductile manner where voids form and grow plastically at large inclusions, which act as hard particles in a plastically deforming matrix. In a recent study of ductile rupture at room temperature of a ferritic steel with 9% chromium [14], voids form a MezlC, particles. Void formation was not observed until strains that reach and exceed the rupture strain of the
on cavitated 12% Cr-Mo-V steel after 5% creep deformation with an energy window around the sulfur KLL energy).
(AES sulfur
56
EGGELER et al.: MICROSTRUCTURAL STUDY OF CREEP RUPTURE
12% Cr-MO-V steel investigated in this work. For the high temperature experiments conducted in this investigation, large oxide and sulfide inclusions (5540 pm, Fig. 3) are more effective in the initiation of voids which grow by plastic deformation than small Me,,C, particles [0.1-0.3flm, Fig. l(a)]. The nature of plastic void formation processes at hard particles does not seem to change at high temperatures for high stresses. It is concluded that with increasing temperatures smaller particles become less effective in plastic void formation and do not play a role in the ductile high temperature rupture of 12% Cr-MO-V steel. 4.2. Cavity nucleation At the low stress level, diffusive creep cavitation appears to be the dominant damage mechanism which results in much lower strains to rupture. Cavities are observed as early as 1% strain [Fig. 6(b)] during primary creep where the strain rate is still decreasing towards the minimum creep rate [Fig. 2(b)]. Middleton [15] observed cavitation in iron based alloys as early as 0.05% strain. It is probable, therefore, that an incubation strain for nucleation is either small or not required for the cavities that form first. Nucleation of cavities occurs continuously during creep and the cavity density is proportional to strain [Fig. 10(b)]; this is also true for other engineering materials [16, 171. If nucleation theory by vacancy condensation is basically correct, it must be concluded that continuous cavity nucleation is a consequence of some stress concentration process which occurs throughout creep life and which triggers cavity nucleation. Argon et al. [27] propose that the triggering event for cavity nucleation might come from the spasmodic nature of grain boundary sliding. It appears unlikely that grain boundary sliding is strongly involved in cavity nucleation for the 12% Cr-M-V steel investigated in this work, since the density of Me,,C, carbides on FAGBs is very high [Fig. l(a)]. The absence of extensive grain boundary sliding is also indicated by the fact that triple point cavities were never found to be wedgecrack like, as one would expect them to be if they were merely a result of grain boundary sliding. Moreover this model cannot explain how cavities are nucleated on transverse FAGBs which do not slide. A more likely mechanism without grain boundary sliding has been proposed by Dyson [16]. It involves the intersection of slip bands with grain boundaries. This mechanism was observed by Nieh and Nix [19] for copper and by Shiozawa and Weertman [29] for a nickle base superalloy, where they found an almost perfect one-to-one correlation between the cavity spacing and the spacing between coarse slip bands. During creep of 12% Cr-Mo-V steel, slip bands do not form [9]. Thus, it is unlikely that cavity nucleation in 12% Cr-Mo-V steel occurs according to the model proposed by Dyson [16].
The AES study conducted in this work suggest that sulphur facilitates creep cavitation in the 12% Cr-Mo-V steel. This interpretation is supported by recent findings of George et al. [ 181. They reported, that for sulfur concentrations above the solubility limit (> 30 ppm in their study), sulphur can segregate between a particle and the surrounding matrix. Once there, it weakens the interface and facilitates cavity nucleation. SEM micrographs like that in Fig. 7 in addition to the Auger results suggest that sulphur segregates at interfaces between the Me,,C, carbides and matrix. This segregation process is certainly dependent on the heat treatment of the as-received material and might also continue during the creep process itself. Assuming that for a given particle size there is a critical concentration of sulphur at a particle/matrix interface which has to be achieved for cavity nucleation, continuous cavity nucleation might also be triggered by continuous segregation. It is also likely that sulphur segregation is only one prerequisite for cavity nucleation, while the nucleation process only starts if other conditions are met as well. Thus the nucleation process also depends on the crystallographic details of grain boundaries [22,23], which can explain why not all boundaries contain cavities. Moreover, if cavitation occurs nonuniformly in the solid, local multiaxial stress states can develop even in an uniaxial creep experiment. This local multiaxial stress states also can influence cavity nucleation and growth during creep [20,21]. 4.3. Location of cavities in the microstructure
In this study cavities are not only found on FAGBs, but also in the interior of the former austenite grains. This is due to the existence of other internal surfaces which form during tempering of the martensitic structure. These are high angle ferrite boundaries and subgrain boundaries [Fig. l(a)], which can also act as segregation sites for sulphur and facilitate diffusion. Thus, cavities in the interior of the former austenite grains do not have to grow by power law creep alone, which might be expected for cavities growing within a grain [24], when no such internal surfaces are present. Since these boundaries intersect former austenite grain boundaries, they also provide transverse diffusion paths for cavity growth on longitudinal FAGBs, where a relative high number of cavities with appreciable size is observed. Although the applied stress would not drive diffusion along a longitudinal FAGB, the growth of such cavities could be facilitated by “atom plating” on intersecting ferrite and subgrain boundaries which are transverse to the stress axis. 4.4. Cavity growth and accumulation of creep damage Although the mean cavity size, evaluated from large specimen regions, increases linearly with creep strain, one should keep in mind that cavitation occurs
jnhomogeReously in the structure. There are sampling fields in Fig. 8 with high and hi cavity densities. Onfy one of the six sam$ing fields wontons two microcracks. Cavitation is a&o i~~omogen~us in the sense that onfy smm of the FAGi% perpendicufar to the stress axis ~~~t~~n cavities. In Rg. 5fb), for example, a FAGB segment near the fracture surface is shown that did not cavitate. It folfows from tbc quantitative ~va~~~tjon preened in Fig. $2, that cavities on FAGBs have a more ~rn~o~ta~t role in the rupture process. They are significantly larger than the caviiies within the former anstenite grtins and are arranged on FAGB facets. The @-parameter @p.mntifying cavity line up along FAG& in this study) is highest and cavity growth is fas@st on FAGB seg~ men@ which are nearly ~er~nd~~~~a~ to the a~~~~~d stress (Fig. 14).
ary facets which are transverse to the apptied stress.
Once a criticat number aE microcracks have coaksced, the situating is no longer one of hon~o~e* neous creep damage but one of creep crack grow&, and the stram rate ia the tertiary crefzp range is strmgiy influenced by the creep damage, The transition from secondary to tertiary creep strongly de* pen& on the statistical positioning 5i” the micros cracks, which are ~etero~eneons~y distribute fRg. 8). This explains why the time to rupture can deviate by a factor of two for the same ~x~e~~me~lta~conditions (see Tabie 2).
A mode1 of the constrained cavity growth process, deposed by Rice [30], is shown sc~~~tj~a~i~ in Fig, li”. A cavitating grain boundary facet ~~i~rneter d) is embedded in creeping body under an apphed stress, 6. The cavities a8 have a radius equal to c and D untform spacing, a. In an extension of the o~~j~~ treatment I3ff, the equation for the cavity growth rate is given as
fn a&r $0 ~~rn~r~ the afo~~~~t~oned wsu&s with theoreticd predictions, we will now study the applicability of the constrained eavity grawth theory, first ~~~~du~ed by Dysan [Zgj. As shown in Fig, 13, ca~~~~t~~~g facets in KY%@r-Me-V steel we t;wpicafIy isolated with respect to other cavitating gram boundaries. Similar&, the constrained cz~&y growth mode1 describes a rn~~~ost~Gtnr~ where isolated transverse grain boundary facets cavitate while the neighboring Kn equation (31, !P is the tip angie of the cavities bo~~~~~~s are und~m~~~. In this easi;r, the cavity d~~~~d by CQSJI= ~~~~~~where yb is the grain bound* growth rate wilt he limited by dislocation creep in the ary free. energy anld ySis the surface fry%energy. The surrounding matrix. sit&ring stress, a,, in equation (1) is equaf to ~~servat~o~~s of cavities early in the He of speci- @y, sin@)jZ_ Vames of the mate~a~ constants used to mens suggest that some of the iasger grain ~und~~~ determine the cavity growth rate using equation {t) cavities nucfeate upon or shortly after loading. As a are given in Table 3. Equation (I} describes thli first ~~~ro~irnat~on, thsr preferred cav~t~t~~~ on trans- growth of qu~i~e~~i~~b~urn cavities as opposed to XW-N grain boundary facets corresponds to the largest flatal crack-like cavities, ff the surface dit%sion is cavities observed in the 12% Cr-Mo-V steef investisu%cientIy SIOW,then the morphotogy of the cavities &a&d here. It fohews that the f~rrn~t~o~ of the first wil$ be Bat arzd crack-Eke, OW ob~er~~~io~~ of cmmicrooraeks could involve these largest cavities, whieb ities in 12% Cr-MeV st&, e.g. Fig. 7, indicais_?that nucleate shortly after loading and be os1grain boundcavities are not crack-like aad typicalXy grow in 8
EGGELER
58
et al.:
MICROSTRUCTURAL
Table 3. Physical constants n
B
(MPa)-66
used in the constrained
Power law stress exponent (650°C)
9.12 x 1O-2’
Power law creep coefficient (650°C)
(m’)
I.18 x 10-29”~21 Atomic
SD,
(m) SF’)
,,I x lO-‘W?
Pre-exponential for grain boundary diffusion
Qb
(J n-m-‘)
1.74 x 105”.*1
Activation energy for grain boundary diffusion
Ys
(J m-*)
2.1’2’
Surface free energy
Yb
(J mW2)
0.8512]
Grain
Plasticityand Creepof
quasi-equilibrium manner in this material. From microstructural measurements on cavitated facets such as the one shown in Fig. 13, a mean value of 20pm was determined for d and mean values of 1 ranging from 5 to 7 pm were determined for grain boundary facets nearly transverse to the applied stress. To a good approximation, a cavity may nucleate and grow if the cavity radius is greater than 2y,/a. This value for 12% Cr-Mo-V steel under 80 MPa, 2yJa = 0.07 pm, was taken to be the initial cavity radius, c,,, in the model. Cavity size is calculated in the model by numerically integrating equation (1). For a comparison with the measured cavity size data, the cavity radius was calculated for an applied stress of 80 MPa and creep times corresponding to 1, 5 and 12% strain. The model predictions are presented in Table 4 where they are compared with the largest 5% of the cavity population. These measured values are used as an indication of the upper limit of the observed cavity sizes. The model successfully predicts cavity sizes for 1 = 5 pm which lie within the measured 5% range. The results for 1 = 7 pm are also good approximations of the maximum observed cavity size demonstrating that the cavity spacing need not be extremely accurate to obtain reasonable predictions. It is evident from these findings that the contained cavity growth model is an accurate description of the grain boundary cavitation in 12% Cr-MeV steel. An unconstrained cavity growth model [32], in which cavity growth is only limited by the grain boundary diffusion, predicts cavity coalescence (2c = 1 = 7 pm) after only 1.3 x lo5 s for the same material and conditions. Since cavities are required to
I
5 I2
Time (s) 8.6 x 10’ 105 1.0 x 106
7.0 x
of upper limit cavity size measurements model calculations
0.6 to 0.9 I.8 1.3 to 1.9
1.2 to
boundary
free energy
Deformation-Mechanism Maps, The Metalsand Ceramics,pp. 62-63. Pergamon Press,
(1987).
Measured cavity radius (rm) [largest 5%]
volume
and M. F. Ashby,
Oxford (1982). ‘21H Riedel , Fracture at High Temperatures,
Strain (%)
cavity growth model
6.6
R
[“H. J. Edward
Table 4. Comparison
STUDY OF CREEP RUPTURE
Calculated cavity radius (rm) 1=5flm 1=7pm 0.7 1.5 1.7
0.8 I.8 2.1
with
Appendix
A. Springer,
Berlin
be present on all facets, cavity coalescence should correspond to the time to rupture for unconstrained cavity growth. The corresponding observed rupture time (approx. 1 x 106s), however, is about an order of magnitude greater than 1.3 x lo5 s. 4.6. Constrained formation
cavity
growth
and
microcrack
There is an important implication of the constrained cavity growth model which has to do with the nature of the stresses acting on a cavitating facet. The facet stress, cr,,in Fig. 17, rapidly decreases early in the course of constrained cavity growth. In this way, the cavitating facet behaves mechanically similar to a traction-free microcrack long before cavity coalescence occurs. Thus, the time to cavity coalescence on isolated facets is not a good criterion for predicting the time to rupture. This is further supported by the fact that, although there is good agreement between the measured and the calculated cavity size values, the calculated time to cavity coalescence (2.8 x lo6 s for 1 = 5 pm) is greater than the observed rupture time by about a factor of 3. Closely spaced microcracks, where cavity coalescence had occurred, were observed late in the life of the specimen. Although cavity coalescence is not predicted for the average grain facet size of 20 pm, it could occur by constrained cavity growth on abnormally large grain facets since the cavity growth rate increases with increasing facet size, d, in equation (1). It is likely that, between two closely spaced cavitating facets, transgranular deformation and damage is accelerated such that the two cavitating boundaries behave as one late facet. In this way the effective facet diameter could be much larger than the grain size. The formation of a large microcrack by this process is shown schematically in Fig. 18. Two closely spaced microcracks can be seen in one of the sampling fields in Fig. 8 which together comprise an effective facet size of about 70 pm. When the value of d in the constrained cavity growth model is 70pm, cavity coalescence (c = J.) occurs earlier than the corresponding observed rupture time. The presence of
EGGELER et al.: MICROSTRUCTURAL STUDY OF CREEP Cavitating Facet Interaction
\
o
I
\
I
Fig. 18. Sequential schematic drawings of a possible mechanism for the formation of microcracks by constrained cavity growth on interacting grain boundary facets.
RUPTURE
59
gation can occur preferentially and which will facilitate diffusion. A quantitative metallographic evaluation of the cavity population showed that cavities on transverse former austenite grain boundaries play the dominant role in the rupture process. The constrained cavity growth model reasonably agrees with the observed growth of closely spaced cavities on “isolated FAGBs”, however, overestimates the rupture time when cavity coalescence is used as a rupture criterion. Therefore, cavity coalescence should not be used as a criterion for creep rupture for a 12% Cr-Mo-V steel. The rupture time is determined by the formation and growth of microcracks from closely spaced cavitating grain boundary facets. Acknowledgements-Financial support of the Swiss Federal Government within the framework of the European cooperation COST 505 is gratefully acknowledged. Joint support from the Swiss Commission for the Promotion of Scientific Research and Sulzer AG, represented by Dr P. P. Schepp and Dr B. Walser, under contract KWF 1512 is appreciated. The authors wish to thank Dr H. Mathieu from the Laboratoire de Metallurgic Chimique for kindly making available the Auger spectrometer and for his help in performing the measurements and Professor W. D. Nix for fruitful discussions during his stay in Lausanne on sabbatical leave from Stanford University.
REFERENCES closely spaced microcracks is therefore consistent with constrained cavity growth theory since the effective facet size is sufficiently large. Although it is not typical for a cavitating facet to interact with another, it appears that the formation and growth of microcracks from the facets which do interact have an important role in determining the life of the specimen. 5. SUMMARY AND CONCLUSION High temperature rupture of a 12% Cr-Mo-V steel was studied for two stress levels at 923 K. At the high stress level (175 MPa) rupture occurs in a ductile manner. Voids form at oxide and sulfide inclusions which act as hard obstacles in a plastically deforming matrix. At the low stress level (80 MPa) rupture is controlled by the nucleation and diffusional growth of creep cavities, the first of which nucleate upon loading. Auger results suggest that segregation of sulfur at a carbide/matrix interface facilitates cavity nucleation which also is dependent on grain boundary crystallography and local state of stress. A characteristic feature of creep cavitation in high chromium ferritic steels is that cavities are not only found on former austenite grain boundaries (FAGBs) but also in the interior of former austenite grains. There are additional internal surfaces such as subgrain boundaries and high angle ferrite boundaries, which have formed during tempering, where segre-
1. J. Z. Briggs and T. D. Parker, The Super 12% Cr Steels. Climax Molvbdenum. New York (1965). 2. B. Walser, d. Brezina’and T. Greiger, krch. Eisenhtitt. 50, 249 (1979). 3. F. Masuyama and H. Haneda, The First Int. Conf on Improved Coal-Fired Power Plants (Proc. Conf), Palo
Alto, California, EPRI (1986). 4. B. Walser Heat and Mass Transfer in Metallurgical Svstems (Proc. Con/X. Dubrovnik. Yueoslavia. 1979.
Hemisphere, Washington, D.C. (1981). 5. R. Sandstrom, S. Karlson and S. Modin, High Temp. Tech. 3, 71 (1985). 6. E. A. Little, D. R. Harries, F. B. Pickering and S. R. Keown, Metals Tech. 4, 205 (1977). 7. I. M. Park, T. Fujita and K. Asakura, Trans. Iron Steel Inst. Japan 20, 99 (1980). 8. J. W. Schinkel, P. L. F. Radermakers, B. R. Drenth and C. P. Scheepens, J. Hear Treat. 3, 237 (1984). 9. G. Eggeler, N. Nilsvang and B. Ilschner, Sreel Res. 58,
97 (1987). 10. G. Eggeler, B. Ilschner, P. P. Schepp and R. Zohner, Mater. Technik. 14, 187 (1986).
I I. H. Weber, Festigkeits- und Bruchverhalten bei hiiheren Temaeraluren. Vol. 2. DV. l-69. Verlag Stahleisen. Dusseldorf, F’. R. G. (1980). 12. B. J. Cane, Mechanical Behauiour of Materials-III. (Proc. Conf.), Cambridge, England. Vol. 2. pp. 1733182, 1979 Pergamon Press, Oxford (1980). 13. N. Nilsvang and G. Eggeler, Prukf. Mefallogr. 24, 321 (1987). 14. B. A. Senior, F. W. Noble and B. L. Eyre, Acra metall. 34, 1321 (1986). 15. C. J. Middleton, Metal Sci. 15, 154 (1981). 16. B. F. Dyson, Scripfa melall. 17, 31 (1983). 17. H. Riedel, Fracrure at High Temperatures. Springer,
Berlin (1987). 18. E. P. George, P. L. Li and D. P. Pope, Creep and
EGGELER
60
et al.: MICROSTRUCTURAL
Fracture of Engineering Materials and Structures (Proc. Conf.), Swansea, England. pp. 169-182. Inst. of Metals,
London (1987). 19. T. G. Nieh and W. D. Nix, Scripta metall. 14, 365 (1980). 20. D. R. Hayhurst and F. A. Leckie, Mechanical Behuviour of Materials-IV. (Suppl). Stockholm, Sweden, (1983).
Pergamon Press, Oxford (I 983). 21. B. J. Cane, Advances in Fracture Research (Fracture 81), 1981, Vol. 3, (Proc. Co&), Cannes, France,
pp. 128551293. Pergamon Press, Oxford (1982). 22. J. Don and S. Mahjumdar, Acta metall. 34, 961 (1986). 23. T. Watanabe, Creep and Fracture of Engineering Materials and Structures, (Proc. Conf.), Swansea, England, 1987. pp. 1555168. Inst. of Metals, London (1987). 24. A. C. F. Cocks and M. F. Ashby, Scripta metall. 16,109 (1982).
STUDY OF CREEP RUPTURE
25. D. J. Gooch, Metal1 Sci. 16, 79 (1982). 26. M. F. Ashby, Fracture 1977, Vol. 1, pp. 1-14 (1977). 21. A. S. Argon, I-W. Chen and C. W. Lau. Creep-Fatigue-Environment Interactions, pp. 4685.
Am. Inst. Min. Engrs, New York (1980). 28. B. F. Dyson, Metals Sci. 10,349 (1976). 29. K. Shiozawa and J. R. Wertmann, Acta metall. 31, 993 (1983). 30. J. R. Rice, Acta metall. 29, 675 (1981). 31. H. Riedel, Mechanical Behaviour ofMaterials-IV, Vol. 1, pp. 67-78. Pergamon Press, Oxford (1985). 32. I.-W. Chen and A. S. Argon, Acta metall. 29, 1759 (1981). 33. For example on page 76 of: Metals Handbook, Vol. 9, Metallography and Microstructures. Am. Sot. Metals,
Metals Park, Ohio (1985).