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Informes Técnicos Ciemat
Corrosion of High Chromium Fenitic1 Martensitic Steels in High Temperahire Water, A Literature Review
P. Fernández J. Lapeña F. Blazquez
925 abril, 2000
Departamento de Fisión Nuclear
Toda correspondent en relación con este trabajo debe dirigirse al Servicio de
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Las solicitudes de ejemplares deben dirigirse a este mismo Servicio.
Los descriptores se han seleccionado del Thesauro del DOE para describir las materias
que contiene este informe con vistas a su recuperación. La catalogación se ha hecho
utilizando el documento DOE/TIC-4602 (Rev. 1) Descriptive Cataloguing On-Line, y la
clasificación de acuerdo con el documento DOE/TIC.4584-R7 Subject Categories and Scope
publicados por el Office of Scientific and Technical Information del Departamento de Energía
de los Estdos Unidos.
Se autoriza la reproducción de los resúmenes analíticos que aparecen en esta
publicación.
Depósito Legal: M -14226-1995ISSN: 1135-9420ÑIPO: 238-00-002-0
Editorial CIEMAT
CLASIFICACIÓN DOE Y DESCRIPTORES
S36
CORROSION; CHROMIUM STEELS; FERRITIC STEELS; MARTENSITIC STEELS;TEMPERATURE RANGE 0400-1000 K STRESS CORROSION; CRACKING
Corrosion of High Chromium Ferritic/Martensitic Steelsin High Temperature Water. A Literature Review
Fernández, P.; Lapeña, J.; Blázquez, F.
64 pp. 35 fig. 40 refs.
Abstract:
Available literature concerning corrosion of high-chromium ferritic/martensitic steels in high temperaturewater has been reviewed. The subjects considered are general corrosion, effect of irradiation on corrosion,stress corrosion cracking (SCC) and irradiation-assisted stress corrosion cracking (IASCC).In additionsome investigations about radiation induced segregation (RIS) are shown in order to know the compositionalchanges at grain boundaries of these alloys and their influence on corrosion properties.
The data on general corrosion indicate moderate corrosion rates in high temperature water up to 350°C.Considerably larger corrosion rates were observed under neutron irradiation. The works concerning tothe behaviour of these alloys to stress corrosion cracking seem to conclude that in these materials isnecessary to optimize the temper temperature and to carry out the post-weld heat treatments properly inorder to avoid stress corrosion cracking.
Corrosion en Agua a Alta Tempertaura de los Aceros Ferríticos/Martensíticos deAlto Contenido en Cromo. Revisión Bibliográfica
Fernández, P.; Lapeña, J.; Blázquez, F.
64 pp. 35 fig. 40 refs.
Resumen:
El trabajo que se presenta en este informe recoge una revisión bibliográfica del comportamiento frente a lacorrosión en agua a alta temperatura de los aceros ferríticos/martensíticos de alto contenido en cromo. Elestudio se ha basado principalmente en valorar la respuesta de este tipo de materiales en corrosión gene-ralizada, efectos de la irradiación en corrosión, corrosión-bajo tensión y corrosión bajo tensión asistidapor irradiación. En esta informe también se recogen algunos estudios de segregación inducida por irradia-ción, con el objetivo de conocer los cambios composicionales que se pruducen en los límites de grano deestas aleaciones y la influencia que estos cambios pueden producir en sus propiedades de corrosión.
Los datos de corrosión generalizada en agua a alta temperatura muestran velocidades de corrosión mode-radas, siendo estas mayores en condiciones de irradiación neutrónica. Los trabajos referentes al compor-tamiento de estos aceros a experimentar corrosión bajo tensión, parecen ser concluyentes de la importan-cia en este tipo de aleaciones de optimizar la temperatura de revenido y realizar apropiadamente lostramientos térmicos posteriores a la soldadura para prevenir este tipo de corrosión.
INDEX
Page
1.-INTRODUCTION 1
2.- GENERAL CORROSION
2.1.- Oxide Layer 2
2.2.- Corrosion Rates 4
3.-EFFECT OF IRRADIATION ON CORROSION 13
4.- STRESS CORROSION CRACKING (SCC) 15
5.- IRRADIATION-ASSISTED STRESS CORROSION CRACKING (IASCC) 25
6.-RADIATION-INDUCED SEGREGATION (RIS) 26
7.- SUMMARY AND CONCLUSIONS 30
8.-REFERENCES
9.-LIST OF FIGURES 37
10-FIGURES 41
1.- INTRODUCTION
A fully martensitic steel containing 10-11% Cr and additions of approximately 0,6% Mo. 0,65%
Ni, 0,25% V and 0,15% Nb, referred to as MANET, has been evaluated for fusion applications in
Europe for first wall and breeder structural component in NET and DEMO0"1. The studies in the USA
on this class of material have concentrated on HT-9 (12% Cr, 0,6% Ni, 1% Mo, 0,3%V and 0,5% W)
and modified 9%Cr -Mo (9%Cr, 0,94%Mo. 0,19 V and 0,18%Nb) steels.
The development of elements-tailored reduced activation ferritic and martensitic steels with 2-
12% Cr and additions of W, V, Ta, Mn, Ti and /or N have also been pursued as part of the US, Japanese
and European Fusion materials programmes^' '. The reduced activation higher chromium martensitic
steels have comparable or higher static and fatigue strenghts at elevated temperatures and superior
fracture toughness (lower ductile-brittle temperatures and higher uppershelf energies in Charpy V-notch
impact test) in the unirradiated condition compared to the conventional -12% Cr-Mo-V-Nb or W steels'4'.
These high chromium martensitic steels have many advantages over the austenitic steels,
including lower decay heat, increased resistances to thermal stress development and irradiation-induced
high temperature (helium) embrittlment and void swelling01. In addition, investigations of the corrosion
of commercial and reduced activation martensitic steels in static and flowing high-temperature, high
pressure water have not revealed any unexpected or abnormal behaviour( ', the corrosion resistance
increases with increasing Cr content, the weight losses for 8-9% Cr steels being approximately twice
those of the 12% Cr steels<4).
There is very little information available concerning the corrosion resistance of low activation
martensitic steels. The following literature review covers the different corrosion phenomena experienced
by martensitic steels in Liquid Metal Fast Breeder and Light Water Reactors.
Pertinent literature concerning general corrosion, stress corrosion cracking (SCC) and irradiation-
assisted stress corrosion cracking (IASCC) of high chromium martensitic steels are reviewed in the
present report.
2.- GENERAL CORROSION.
High temperature corrosion is a complex mixture of interrelated processes and reactions. After the
initial nucleation and the formation of a continuous film or reaction products scale has been formed, the
metal and the reactant are separated and the reaction proceeds through diffusional transport of reactant
atoms or ions through the scale. In certain cases electron transport through the scale may alternatively be
rate-determining process.
In addition to these transport processes other phenomena occur. Thus grain growth occurs in the
reaction products, and if more than one reaction product is formed, these may react to form new
corrosion products. Furthermore, growth stresses are built up in the scales, and these are alleviated
through various mechanisms which may include high temperature creep and deformation, cracking and
/or spallation. Microchannnels are also probably formed through these processes. It is also common that
porosity and voids develop in the scales through the deformation processes and as a result of scale
growth by outward transport of metal ions through the scale.
2.1 Oxide Layers:
General corrosion in water and steam on Ferritic/Martensitic steels have been estudied extensively
in steam generators, superheaters of the Liquid Metal Fast Breeder Reactor (LMFBR). Generally, in
these types of materials the oxide layer formed is magnetite. The investigations realized in this area were
focused with the objetive to establish oxide morphology, corrosion kinetics, metal loss and long-term
data concerning oxide adherence in function of water chemistry and heat flux.
Aqueous corrosion of high chromium ferritic or martensitic steels results in the formation of
double oxide layers'3'6'7'. Tomlinson et al(7'8), studied the deposition of magnetite on chromium ferritic
steels (Table 1) in high temperature water (350°C) at high heat flux (0 to 860Kw/m") using a sodium
heated test section. In most tests, all volatile treatment (AVT) water chemistry was used (pH 8,5-9,2 at
25°C controlled by ammonia addition, dissolved oxygen < 7 ppb, Fe < 10 ppb, total dissolved solids < 50
ppb, Cl < 10 ppb, Si O2 < 2 ppb, Cu < 2 ppb). Some tests were also conducted under water chemistry
fault conditions, with additions of sodium hydroxide, sodium bisulphate and oxygen.
Alloy
2,25Cr-lMo
9Cr-lMo
9Cr-lMo
C
0,10
0,10
0,08
Cr
2.27
9,14
8,79
Mo
0,95
1,05
0,95
Nb
0.90
-
-
Si
0,35
0,76
0,61
Ni
0,64
0,23
0,27
Mn
0,50
0.51
0,52
S
0,005
0,002
0,006
Cu
0,02
0,01
0,02
Table 1: Chemical composition of the alloys investigated.
The outer layer consists of magnetite crystals of well developed tetrahedral or octahedral
shapes with diameters > ljim. The magnetite crystals are precipitated from solution and show an epitaxy
with underlying metal surface. Under certain conditions the outer layer crystals coalesce and form a
protective layer on the surface, see Figure 1,3b. Under other conditions, for example when precipitation
of magnetite can occur on other surfaces in the system, the outer layer can be non-protective, see figure
l,3a,4a. The inner layer grows inwards from the original metal surface. It consists in general of a fine-
grained chromium - iron spinel structure and controls the rate of corrosion '. The inner layer has in
general the same thickness as the consumed steel.
Figure 1: Schematic development of double oxide layer.
2.2.- Corrosion rates
Several investigations of general corrosion of high-chromium ferritic steels have been performed
with regard to applications in thermal and fast reactor steam generators. The general corrosion behavior
of 2l/4Cr-lMo steel has been studied in aqueous solutions, in pure and impure superheated steam, in
saturated steam, and under nucleate boiling conditions, with chloride and oxygen additions' ' '.
Oxidation is typically measured either as metal weight loss, as an oxide film thickness, or wall
penetration depth. Oxidation is a function of solution pH, increasing dramatically at low or very high
values.
Under laboratory (simulating the chemistry of LMFBR steam generator) conditions a relatively
stable protective oxide layer is formed; extrapolation to 30 years plant lifetimes gives 100 to 130 ¡¿m
metal consumption in the temperature regime 497°C-527°C. However, recommended allowances are
much higher (510-760 ¡j,m) to allow for periodic loss of scale integrity and other causes( }.
Tomlinson et al ( ' } in a study of magnetite deposition also determined the corrosion rates of 2
1/4% and 9%Cr-lMo ferritic steels (Table 1) in a pressurized water loop at 350°C. The majority of the
test were carried out in all volatile treatment water chemistry, but some tests were also conducted under
water chemistry fault conditions, with additions of sodium hydroxide (9.0 ppm), sodium bisulphate (2.0
ppm) and oxygen (56 ppb).
The variation of metal loss with time is shown in figure 2. Within the wide scatter band of the
data, the rate of corrosion appears to be independent of steel type of (21/4% Cr or 9%Cr). The
experimental data also show (within the same range of scatter) that the corrosion rate is independent of
water chemistry (except where a high level of dissolved oxygen and heat flux are present at the same
time), and thickness of deposited magnetite (in the range 0-7jjjn). A least squares fit of the data gives the
following expression for metal loss versus time:
m = 0,23t0j9"' where, m = metal loss in jam and t = time in days.
The oxide deposited throughout the test section was normally magnetite. However, under one particular
set of conditions (high dissolved oxygen level together with heat flux) a-Fe:O:, was deposited. The oxide
deposited in non-heat flux regions of the same test section was almost entirely Fe O-t. The effect of these
conditions upon metal loss is shown in Figure 3, together with the results from a typical "All volatile
treatment" (AVT) water chemistry run. As can be seen, metal loss is lower (by 27%) in the region of a-
Fe:C>3 deposition compared to elsewhere.
In other corrosion studies related to steam generators the environments used in general have been
water with chloride ions. Broomfield et al"'* investigated the corrosion of a number of steels, including
AISI type 405 and 410, in O,1M NiCb solutions at 300°C. After an initially steep increase the thickness
of the inner oxide layer was found to increase linearly with time . Mean results for the type 405 and 410
steel specimens were approximately 45u,m at 1200 hours.
Corrosion of a few high chromium steels in approximately 300°C deoxygenated water containing
various concentrations of metal chlorides, was also studied by Vaia et al( J>. The composition of the
various 12%Cr stainless steels used in these investigations are presented in table 2.
Alloy
405
405
405
410
409
Heat
2804-1
1781-1
1782-2A
2926-1
2725-1B
C
0,064
0,054
0,056
0,070
0,039
Mn
0,51
0,53
0,53
0.55
0,41
P
0,015
0,014
0,017
0,015
0,023
S
0,025
0,004
0,006
0,008
0,004
Si
0,66
0,56
0,55
0,44
0,60
Cr
13,40
13,55
13,40
12,30
10,83
Al
0,100
0,130
0,110
-
-
Ni
0,12
0,41
0,36
0,22
0,32
Ti
-
-
-
-
0,280
Table 2: Composition of tube support plate materials.
In the first series of tests, the corrosion rates and corrosion products were characterized after
exposure to concentrated solutions. The second series of tests, in concentrated solutions, used a creviced
specimen (isothermal capsule) to compare the relative rate of growth of corrosion product thickness.
Finally, a more realistic assembly (heated creviced assembly) was tested to simulate the concentration
processes ocurring in the crevice region.
Corrosion of the 12%Cr stainless steel in deoxygenated chloride solutions occurs by dissolution
and solution transport of iron ions away from corroding surface. The resulting deposit is characterized by
a double layered oxide. The inner layer is compact, regular and apparently protective while the outer
layer is made up of loosely adherent large tetrahedral crystals. The oxide found on the 12%Cr stainless
steel occupied the same volume as the original metal. This results agrees favorably with work performed
by Mann and Teare( \ which showed that the inner layer oxide formed on an 8%Cr-Fe alloy, after
exposure to a O,1M NiCb at 300°C, occupied 95% to 100% of the volume of the virgin metal oxidized.
The inner oxide layer found on the 12% Cr stainless steel specimens may be the result of partial
dissolution or recrystalization of a barrier layer at the metal surface. The inner layered oxide observed on
the 12% Cr stainless steel has been shown to be Cr rich. The increased Cr concentration within the inner
oxide layer is related to the low solubility of Cr ions in deoxygenated aqueous chloride solutions. It is the
formation of this Cr rich inner oxide layer which imparts the improved corrosion resistance for the 12%
Cr stainless steel. Any factors which may increase the solubility of Cr in the aqueous chloride
environment would lead to reduced corrosion resistance of Cr stainless steel. During the initial stages of
corrosion, Fe ions migrate to the surface, at which, they form the outer layer by precipitation. The Cr
ions, having lower solubility in the bulk chloride environment, tend to form an in situ Cr rich oxide. As
the corrosion continues, the rate controlling step becomes the solution transport of Fe ions across the Cr
rich oxide. It has not been well established whether the controlling step is related to Fe ion transport
across the inner oxide layer or to Fe ion transport across a thin barrier layer adjacent to the metal surface.
Corrosion tests of developmental low-activation martensitic stainless steel were performed by
Ashmore and Large(15>. The behaviour of the investigated alloys (LA7Ta, LA12Ta and LA12TaLC), was
compared with that of a standard martensitic stainless steel, FV448. In addition, some samples of welded
and unwelded LA12TaLC plate were compared (Table 3).
Alloy
LA7Ta
LA12Ta
LA12TaLC
FV448
C
0,15
0,16
0,09
0,10
Si
0,07
0,03
0,03
0,46
Mn
0,76
0,80
1.01
0,86
Cr
11,4
9,8
8,9
10,7
Ni
0,02
0,02
0,02
0,65
Mo
<0,01
<0,01
<0,01
0,60
Nb
<0,01
<0,01
<0,01
0,26
Ta
0,12
0,10
0,09
V
0,25
0,27
0,39
0,14
W
2,9
0,85
0,76
N
0,07
0,04
0,02
Table 3: Composition of the steels tested.
The two sets of corrosion tests, of 500h and 1000h duration were performed in a static autoclave.
The water chemistry conditions used for corrosion testing were representative of PWR primar}' circuit
coolant at a temperature of about 300°C. High purity water was dosed with Lithium hydroxide (LiOH)
and boric acid (H3BO3) to give 2,2 mg Kg" Li and 1200 mg Kg' B. The water was deoxygenated by
flushing with 5% hydrogen in argon, and an overpressure (about 500 psi) of this gas mixture was
established before the corrosion tests were started. After exposure, all specimens were then weighed to
determine the weight change during corrosion.
The results of measurements made on specimens after testing for 500 hours are given in Table 4.
All samples showed overall weight gains during the test. The average weight gains for the FV448,
Lal2Ta and LATaLC steels were identical, while the average gain for the LA7Ta steel was lower by a
factor of three. Because the specimen weight change is the result of two processes - loss of iron into the
water by corrosion giving a weight loss and deposition of magnetite onto the surface giving a weight
gain- it is impossible to derive corrosion rates from weight change data alone.
However, the significant difference between LA7Ta and the other steels suggests that there is a
difference in corrosion and /or deposition rates. The size of the crystals in the outer (deposited) oxide
layer are very similar for all four steels. However, there appear to be fewer crystals on LA7Ta. This
observation is consistent with the oxide thickness measurements, which show LA7Ta to have slightly
less deposited oxide (Table 4).
7
Steel
FV448
LA7Ta
LA12Ta
LA12TaLC
Weight
change
during the
test (nig)
0.32
(0.12)
n=8
0.11
(0.09)
n=S
0,32
(0.12)
n=S
0.32
(0,11)
n=8
Crystal
size and
range (ujn)
1.3
(0.7)
0.3-4.0
1.5
(1.1)
0.3-7,8
1.2
(0.6)
0.3-3.4
1,3
(0.7)
0.5-6.3
Oxide Thickness by Optical
Microscopy (fini)
Total
oxide
0.98
(0.18)
n=17
0.82
(0.15)
n=20
1.04
(0.22)
n=22
1.30
(0.24)
n=29
Inner
layer
0.64
(0.10)
n=17
0.59
(0.06)
n=20
0.73
(0.10)
n=22
0,84
(0.14)
n=29
Outer
layer
0.34
(0.11)
n=!7
0.23
(0 .11 )
n=20
0.31
(0.17)
n=22
0.46
(0,20)
n=29
Oxide Thickness by
Descaling (ujni)
Total
oxide
0.85
(0,04)
n=3
0.50
(0.02)
n=3
0.95
(0,06)
n=3
0,90
(0.06)
n=3
Inner
layer
0.67
(0.03)
n=3
0.39
(0.04)
n=3
0.81
(0.07)
n=3
0.74
(0,05)
n=3
Outer
layer
0.18
(0,04)
n=3
0.11
(0.02)
n=3
0,14
(0.02)
n=3
0.16
(0,02)
n=3
Depth of
Corrosion
(Uni)*
0.52
80.05)
n=4
0.31
(0.02)
n=3
0.57
(0.04)
n=3
0.54
(0,04)
n=3
Values in brackets represent one standard deviation, n is the number of observations.
* Calculated from the difference between the initial specimen weight and the specimen weight
after descaling, with the density of the steel taken to be 7,41 g cm°.
Table 4: Average values of measurements on disc specimens (500h test)
Results of measurements made on plate specimens after 500 hours of testing are given in Table 5.
The weight changes recorded after corrosion appear very different from those of disc specimens (table
4). The unwelded plate specimens (LA12TaLC 1) show an average weight gain some four times higher
than the equivalent disc specimens. When adjusted to allow for the higher surface area of the plate
specimens (plate specimens about 9 cm", disc specimens about 6 cm") the difference is still significant.
All the welded plate specimens (LATaLC 2 and LA12TaLC 3) showed weight losses (Table 5) with
LATaLC2 having higher weight losses than LA12TaLC 3.
Steel
LA12TaLC 1
(Unwelded)
LA12TaLC 2
(Welded)
LA12TaLC 3
(Welded)
Weight
change
during the
test (mg)
1.34
(0.02)
n=3
-1.61
(0,97)
n=3
-0.58
(0.25)
n=4
Region of
Specimen
Measured
Edge
Edge
Weld
Edge
Weld
Oxide Thickness by Optical
Microscopy (pjm.)
Total
oxide
1.17
(0.12)
n=6
1.57
(0.19)
n=S
1.43
(0.69)
n=16
1.72
(0.24)
n=9
2,20
(1.37)
n=22
Inner
layer
0.77
(0.09)
n=6
0.54
(0.06)
n=8
0.45
(0.08)
n=16
0,67
(0.08)
n=9
1.62
(1,65)
n=22
Outer
layer
0.40
(0.10)
n=6
1,03
(0.16)
n=S
0.97
(0,63)
n=16
1.05
(0,20)
n=9
0.58
(0,54)
n=22
Oxide Tliickness by
Descaling (uni)
Total
oxide
1,45
(0.39)
n=2
1.70
(0.14)
n=2
1.16
(0,12)
n=2
2,63
(0,33)
n=2
1.43
(0,09)
n=2
Inner
layer
0.96
(0.21)
n=2
0.91
(0.05)
n=2
0,76
(0,07)
n=2
0.84
(0.13)
n=2
0.89
(0.11)
n=2
Outer
layer
0.49
(0.18)
n=2
0.80
(0,10)
n=2
0.40
(0,05)
n=2
1.81
(0,22)
n=2
0.54
(0.01)
n=2
Figures in brackets represent one standard deviation, n is the number of observations.
Table 5: Average values of measurements on plate specimens (500h test).
Results of measurements made on disc specimens at 300°C during 1000 hours are given in Table
6. All specimens showed overall weight gains, as in the 500-hour test (Table 4). The average weight
gains for the FV448, LA12Ta and LA12TaLC steels were again identical and were about double the
500-hour test values. The average weight gain for the LA7Ta steel was again lower than the weight gains
for the other steels, but the difference was much less marked.
Steel
FV448
LA7Ta
LA12Ta
LA12TaLC
Weight
change
during the
test (nig)
0.62
(0.19)
n=8
0.47
(0.14)
n=8
0.62
(0.18)
n=8
0.62
(0,18)
n=8
Crystal
size and
range (jim)
2.1
(0.7)
1,1-4.8
2.0
(0.5)
1.0-3.8
2.1
(0.6)
1.1-4.0
0.1
(0.7)
0.6-5.2
Oxide Thickness by Optical
Microscopy (urn)
Total
oxide
1.38
(0.16)
n=20
1,36
(0.25)
n=15
1.54
(0,18)
n=15
1.52
(0.18)
n=16
Inner
layer
0.70
(0.07)
n=20
0.58
(0.11)
n=15
0.63
(0,10)
n=15
0.62
(0,07)
n=16
Outer
layer
0.68
(0.13)
n=20
0.78
(0.20)
n=15
0.91
(0.24)
n=15
0.89
(0.15)
n=16
Oxide Thickness by
Descaling (urn)
Total
oxide
1.28
(0.09)
n=3
0.99
(0.009
n=3
1.30
(0.07)
n=3
1.37
(0.02)
n=3
Inner
layer
1.12
(0.09)
n=3
0.82
(0.05)
n=3
1.11
(0.03)
n=3
1.20
(0.04)
n=3
Outer
layer
0,16
(0.04)
n=3
0.17
(0.05)
n=3
0.19
(0,05)
n=3
0,16
(0.01)
n=3
Depth of
Corrosion
(fini)*
0.75
(0,09)
n=3
0.57
(0.01)
n=3
0.78
(0.02)
n=3
0,80
(0,05)
n=3
Table 6: Average values of measurements on disc specimens (lOOOh test).
The average diameter of the crystals in the outer layer is similar for all four steels, though the
crystals are more numerous and slightly larger than those after the 500-hour test. In the 1000-hour test no
difference is distinguishable between the number of crystals per unit area on LA7Ta and on the other
three steels.
The average values of measurements on plate specimens for 1000 hour of test are shown in Table
7. The pattern of weight changes recorded is similar to that found after the 500 hour test (Table 5), with
the unwelded plate (LA12TaLC 1) specimens showing weight gains and the specimens of welded plate
(LA12TaLC 2 and 3) showing weight losses. Whilst LA12TaLC 3 specimens had a similar average
weight loss to that found after the 500 hour test, LA12TaLC 1 had a lower average weight gain and
LA12TaLC 2 had a lower average weight loss.
10
Steel
LA12TaLC 1
(Unwelded)
LA12TaLC 2
(Welded)
LAllTaLC 3
(Welded)
Weight
change
during the
test (mg)
0,32
(0.08)
n=2
-1.61
(0.97)
n=3
-0,58
(0,25)
n=4
Region of
Specimen
Measured
Edge
Edge
Weld
Edge
Weld
Oxide Thickness by Optical
Microscopy (um)
Total
oxide
1.66
(0.15)
n=10
1.43
(0.25)
n=9
1.77
(0.76)
n=6
1.48
(0.15)
n=8
1.99
(0.93)
n=17
Inner
layer
0.66
(0.08)
n=l0
0.60
(0.06)
n=9
0.65
(0.10)
n=6
0.59
(0,08)
n=8
1.16
(1.25)
n=17
Outer
layer
1,00
(0.16)
n=10
0.83
(0.26)
n=9
1.12
(0,68)
n=6
0.88
(0,17)
n=8
0.83
(0.55)
n=17
Oxide Thickness by
Descaling (uni)
Total
oxide
1.45
(0.28)
n=2
2.15
(0.13)
n=2
1.57
(0,45)
n=2
1.79
(0.04)
n=2
1,27
(0.05)
n=2
Inner
layer
0.92
(0.01)
n=2
1.34
(0.12)
n=2
0.94
(0.27)
n=2
1.28
(0.18)
n=2
0.91
(0.18)
n=2
Outer
layer
0.53
(0.29)
n=2
0.82
(0.01)
n=2
0.56
(0,08)
n=2
0,51
(0.23)
n=2
0,36
(0.13)
n=2
Table 7: Average values of measurements on plate specimens (1000h test)
Estimates of the total oxide thickness and the thickness of the inner and outer layers have been
made by three methods-optical microscopy, chemical descaling of the oxide layers, and determination of
the weight of metal consumed in the oxidation process. The last of these methods leads to an estimate of
the inner layer thickness only.
The measurements made on disc specimens indicate that two of the low activation steels, LA12Ta
and LA12TaLC, exhibit corrosion properties very similar to those of the reference steel, FV448, and
follow an approximately parabolic corrosion rate law. The other steel, LA7Ta, shows a lower corrosion
rate althought the difference becomes less marked as the exposure period increases. It also appears to
show a departure from parabolic kinetics, making prediction of long-term behaviour more difficult.
However, it would be unwise to draw conclusions with regard to corrosion kinetics on the basis of
measurements for only two periods of exposure.
11
The difference in corrosion kinetics of LA7Ta is probably associated with the high level of
tungsten in this steel (2,9 wt%) compared to the other steels (0-0.85 wt%). Whether the tungsten affects
the corrosion process directly, or via the effects of cold working cannot be deduced from the present
data. The question could be resolved by the application of surface analytical techniques to determine
elemental profiles in the specimen surface regions.
Althought the present tests are restricted to a limited set of temperature, pressure and water
chemistry conditions, they indicate no undue sensitivity of the LA steels to aqueous corrosion.
The measurements made on plate specimens only allow the effect of welding on the low
activation steel LA12TaLC to be assessed. Since the corrosion properties of most alloys are modified in
the weld area, and in the associated heat affected zone (HAZ), it is not possible to say how this steel
compares with the reference material or with the other low activation steels in this respect. Interpretation
has been further complicated by the presence of a thick scale and regions of internal oxidation or
contamination at locations away from the weld.
Data from the regions of plate specimens, which could be interpreted (weld material and cut
edges) suggest corrosion behaviour broadly comparable with that found on the disc specimens of
LA12TaLC steel. No evidence was found of intergranular attack or cracking in the weld regions.
In order to make a sound assessment of the corrosion behaviour of welded low activation steels it
would be necessary to compare welded specimens of the low activation steels with equivalent specimens
of the reference steel. In addition it is important that the specimen be free of both oxide scale and
internal oxidation or contamination. These features should be removed prior to welding because the heat
input and cold work produced during the process would be likely to modify the specimen surface
properties.
Uniform corrosion tests of E A Heat F-82H modified low activation martensitic steel (7.65Cr,
2.1W, 0.100C, 0.16Mn, 0.14V, 0.003S, 0.002Ta wt%, balance Fe) have been carried out by Lapeña et
al1 '. In these series of tests the F-82H modified was tested on samples from welded material plates
(TIG and EB).
In all cases specimens were cut from each plate in rectangular form of 2-3 mm thickness, 13-
15 mm width and 50-60 mm length. Some of them were only from base material (normalized at
1040°C/37' plus tempered at 750°C/lh air cooled) and others contained the weld and heat affected
zone (HAZ) in the middle of the specimens. The samples were tested up to 2573 hours in a
12
recirculated autoclave at 260°C in water with 0,27 ppm of lithium and 2 ppm of hydrogen at room
temperature. Different extarctions were performed at various intervals. After samples were removed
from testing and before the final weight, the specimens were electrochemically cleaned with CNNa to
remove any corrosion products. This reactive permits the removal of the oxides without producing
material attack.
Recent results of weight losses and weight loss rate up to 5000 h have been reported by
Lapeña et al (1 '. Results can be seen in figure 4. Weight losses are not very different in all the
materials state tested, although seems a little higher in the EB weld material (see Fig. 4). They are
about 30 mg/drrf after 500 hours, 60 mg/dm~ after 2573 hours and 80 mg/dm" after 5000 hours test.
The weight loss rate for the base metal and weldment material seems to stabilize after 5000 hours to a
value of about 0,01 mg/dm" h. Yamanouchi ( ', in a summary of engineering data for use of reduced
activation martensitic steel, showed weight losses of F-82H (experimental heat) of approximately 162
mg/dm' . However, the results are not comparable, because the testing conditions are different. The
weight losses of 162 mg/dm2 were obtained testing the material in water, without additives, at 250°C
with 200ppb 0a during 250 hours.
3.- EFFECT OF IRRADIATION ON CORROSION
Neutron and y-irradiation are expected to influence the corrosion rates owing to the radiolytic
decomposition of water. In some systems the production of additional oxidizing species is reported to
accelerate corrosion, while in others it may cause a reduction in dissolution rate through the production
of a protective oxide film. One other posibility, with metals which are protected by passive oxide layers,
is that the increase in the metal rest potential resulting from irradiation could lead to the initiation of
i i- J • (19.20)
localized corrosion
Results from an investigation of the effect of irradiation on the corrosion of specimens of various
martensitic steels are reported by Källström'"11 and in part by Gott and Lind'"'. The following alloys and
kinds of specimens were included: 1.4914 (MANET), 1.4914 with electron beam weld and with laser
weld, FV448, LA7Ta, LA12Ta and LA12TaLC (alloy composition given in Table 3). Test coupons were
inserted in stainless steel holders assembled in the three different sets in a high pressure water loop in the
Studvik R2 Reactor. One of the sets of holders was located in the core part of the loop, a second set
above the core where the specimens were exposed to some products of radiolysis but no radiation. The
third set of coupons was placed in the loop away from the core, for reference. The loop water
temperature was 275°C and the water velocities 3,5 ms"1 in the core and about 0,2 ms"1 at the reference
13
specimens. The thermal and fast neutron fluxes were about 1,0. 10 n/m's. Exposure times were
approximately 300. 1500 and 5000 hours. The extent of corrosion was determined from the weight
changes after ultrasonic cleaning and descaling (only ultrasonic cleaning of reference specimens).
The corrosion of reduced activation steels is slightly lower than that of the 1.4914 MANET type
steels tested in the reference region of the loop. The electron beam and laser welding do not appear to
have significantly affected the corrosion of the 1.4914 steels. Whilst the corrosion of the FV448 steel is
comparable to that of the 1.4914 steels following exposure in the reference and radiolysis regions, the
magnitudes of metal consumed by corrosion are significantly lower than those of the 1.4914 and reduced
activation steels tested for 300 hours in the core region.
Corrosion weight losses of FV448 and reduced activation steels tested for 500 and 1000 hours in
high purity water at 300°C and 16MPa pressure in static stainless steel autoclaves by Ashomore and
Large03' are considerably lower than for the same steels exposed in the reference region of the loop in
the present study. The lower corrosion rate of the LA7Ta steel in the Ashomore and Large investigation
was tentatively attributed to its higher tungsten content but this superiority is not clearly apparent in this
work. The major differences in the results of the respective studies are probably due to the effects of the
water flow on the corrosion kinetics.
The weight losses due to corrosion of the coupons exposed in the reference region of the loop
follow at"1" relationship. The weight losses after 5000 hours exposure are in the region of 3mg/cm~~
corresponding to a thickness reduction of approximately 4 ¡xm. The corrosion of steel coupons tested in
the core region is considerably larger than that of the coupons exposed in the radiolysis region of the
loop. This is probably a consequence of the higher concentrations of radiolytic products in the core
region and possibly the reduced chromium contents in the steel matrices due to irradiation induced a'
(Cr-rich ferrite) precipitation. The corrosion is even lower in the reference region where the
concentrations of the radiolytic products are negligible. Contrary to expectation, the weight losses after
descaling are less for coupons with a long exposure time than those with a short exposure, and
particular/ for those exposed in the core region. However, it must be remembered that the coupons have
not been exposed at the same location. The behaviours are tentatively attributed to progressively
increasing crud deposition preventing access of the water to the surfaces and thereby reducing the extent
of the corrosion.
14
4- STRESS CORROSION CRACKING (SCC)
Various definitions of stress corrosion cracking (SCC) have been proposed; the one adopted is the
brittle or quasi-brittle fracture of a material under the conjoint actions of a stress and a corrosive
environment, neither of which would cause such fracture acting alone or consecutively.
Stress corrosion cracking of martensitic and ferritic stainless steels has been extensively reviewed
by J.E. Truman'"3'. Tempering of 12-13% Cr martensitic steels at temperatures in the range 350 to 600°C
was found to be particulary detrimental to the SCC resistance. This paper studies the susceptibility of 12-
13% martensitic stainleess steels from different points of view; a) Effect of strength, carbon content and
heat treatment, b) Effect of applied stress, c) Effect of temperature, d) Effect of polarization, e) Effect of
environment, f) Irradiation experiments, g) Effect of prior exposure to a corrodent, h) Effect of steel
composition, i) Notched specimens, j) Precracked specimens.
Depending on the carbon content and the heat treatment applied, 13% chromium steels can
conform to a low, medium, or high strenght designation and, indeed, some steels can do so very well
simply because of the choice of heat treatment. To give a low-strenght condition, all steels must either
be cooled very slowly from the austenitic temperature range or hardened and then tempered at a
temperature in excess of 650°C. From the data of Figs (5, 6, 7 and 8), and also from many very long-term
tests, it is know that under such conditions resistance to cracking is extremely high. Softened steel has
withstood aggressive conditions in the laboratory without cracking and has been used widely in industry
without stress corrosion being a hazard. It seems probable that these chromium steels passivate much
more readily than do the simple low-strenght steels, and so cracking is less likely under the specific
environment conditions potentially dangerous with the latter.
For the medium and high strenght designation, it is obvious from the data of Fig 9, and 5,10,11
that steels of higher strength are progressively more susceptible if the strenght is achieved by a standard
hardening and tempering sequence, but that the tempering treatment used can also affect resistance.
When environments are sufficiently acid, hydrogen may be accepted as a cathodic reaction
product and, even when the bulk solution is not sufficiently acid for hydrogen evolution, local
acidification by the hydrolysis of metal ions at stagnant points and pits may be proposed. In a given
environment, steels of lower strenght may resist failure for longer periods, or indefinitely, however
failure of the lower -strenght steels could well occur in a more agressive environment, i.e. one which is
more acid and/or contains cathode poisoners, Figs. 5 and 8. The results of polarization tests are in
agreement with an Hydrogen Embritlment (HE) model. In a neutral chloride solution both cathodic and
15
anodic polarization accelerate brittle fracture, the former by cracks which appear to nucleate around the
full circumference of the specimen, and the latter by cracks which grow from isolated points and thus
can be explained in terms of local acidification (Fig 12). With a sulphuric acid solution, however, only
cathodic polarization led to brittle fracture, anodic attack simply caused general corrosion. Anodic
polarization in a sulphuric acid solution causes general anodic dissolution by transpassivity and thus
there is no variation from place to place in corrosion mode or conditions. Moreover, the surface potential
is uniform and above that at which hydrogen could be evolved. Anodic polarization in a near neutral
chloride solution, however, leads to the production of pits with acidification and the possibility that local
environment potential conditions are such that hydrogen may be evolved locally. Increasing the
environment temperature has variable effects on cracking rate (Fig 5 and 6) but it has been shown that
HE is less pronounced at elevated temperatures, althought still possible at temperatures as high as
200°C. As with many SCC data produced using plain specimens, it is possible to obtain a straight-line
relationship by plotting initial stress versus the logarithm of time to rupture, although there can be a
change of slope at some stresses (Fig 13 and 14). The slope may vary according to stresses (Fig 14) or
environment (Figs 13 and 14). For SCC to be the failure mechanism there must, by definition, be a
threshold stress below which cracking is not possible. With the higher strength state in more agressive
environments (Fig 13) this threshold stress must be very low and, althought cracking can be very
delayed, it cannot be assumed that it is other than low in other environments. Most data for 13%
chromium steels have been obtained using smooth specimens, although some results of testing based on
the application of linear elastic-fracture mechanism are available. By comparing the results of Figs. 13
and 15, it can be seen that the stress-concentrating effect of a precrack reduces the time to rupture
drastically, whilst the low values of Kiscc indicate that the threshold stress is probably very low. It is
debatable whether design should be on the basis of a maximum defect size to ensure that Kiscc is never
reached. The limited amount of results of testing using "engineering" notches, listed in table 8, gives a
somewhat variable picture, presumably because coincidence of corrosion initiation with the defect is
necessary for a crack to start and this is less likely in many environments than is the case with non-
stainless steels.
No such variability is apparent for the notched sample, hydrochloric acid test results (Fig. 7) show
depassivation in this medium is rapid and general. It could be argued that by ignoring the initiation
period for cracking on a plain surface one is discarding a potentially valuable property. Certainly, in the
context of stainless steels the casual dismissal of the initiation stage as worth while parameter is
questionable, althought the arguments may be valid for some materials of doubtful engineering value.
The indications are that the initiation resistance is a material property. From a practical point of view, it
should be noted that a portion of the initiation time may be used up by environmental exposure without
stress.
16
Environment at room
Temperature
Atmosphere
3% NaCI
Notched*
Yes
Yes
No
Yes
Yes
No
Time to rupture (h)
Tempered 250°C/2h
(HV 517)
186
310
2580
Time to rupture (h)
Tempered 450°C/2h
(HV 531)
210
7,2
144
2,3
70
100
Notch* : included angle 47,5°, tip radius 0,38 mm. root diameter 3,4 mm.
Table 8: Effect of presence of notches on time to failure of specimens of hardened and tempered
13% Cr steel stressed in presence of corrodent, initial stress 310MNm~~.
From the precracked test data, it has been shown that the crack propagation rate varies with stress
intensity in the same way as for many materials (Fig 16) with regions of marked K dependence (zones I
and HI) at values above and below "plateau" (zone II).
The sodium chloride content of the solution used had little effect on Kiscc (Fig 17) in contrast to
the use acid solutions did (pH 3 and 1,5). This latter may be considered a little surprising since, by the
mechanism of crack solution control proposed, the pH should rise or fall to an equilibrium value.
Possibly limited sample size allowed the penetration of hydrogen ions to the crack tip by diffusion.
Polarization either anodically or cathodically had no effect on the cracking rate in zone U, but reduced
the time to rupture possibly by affecting the crack-propagation initiation time. Potentciostatic
polarization markedly affected the time to rupture (Fig 18). Unpolarized, such steels take a potential of -
-350 mV (SCE) in 3% sodium chloride solution and corroding. Increasing potential will stimulate
corrosion and reducing potential will decrease corrosion, with none below -700mV althought hydrogen
evolution is possible.
All the data discussed so far may be considered explainable in terms of a HE mechanism.
However, if steels tempered at between 350°C and 650°C are considered, there can be evidence of an
Active Path Corrosion (APC) effect. Whether tempered at 250°C, 450°C or 550°C, there is a good
correlation between strenght and the logarithm of time to rupture (Fig. 10 and 11) with, at a given
strength, 250°C tempering giving the most favourable results and 550°C the most adverse. Simple
tempering temperature versus logarithm of time to rupture show 450°C to be the worst tempering
17
temperature for a given steel (Figs. 5 and 6) but this may be attributted to the marked loss in strenght
introduced by tempering at temperatures above 450°C (Fig. 9). Thus a HE effect may still be assumed
with material tempered at temperatures above 450°C, but another feature also appears likely. The
resistance "trough" at 450°C (Figs. 5 and 6) may be associated with loss of toughness, obviously of
relevance to a brittle-fracture mechanism, if only in establishing the amount of cracking required to give
unstable growth (Fig. 9). The effect of tempering on corrosion resistance in the range 450-650°C may
also be of relevance since this is associated with selective corrosion with preferential attack along prior
austenite grain boundaries. Such selective attack is attributed to localized chromium depletion caused by
growing chromium-rich carbides. It should be noted that cracking characteristics of steels ruptured in
sodium chloride solution or the atmosphere vary according to tempering treatment, being markedly
intergranular with material tempered at 450°C or higher, and partially transgranular with steel tempered
at 350°C or below. Since cathodic polarization led to transgranular cracking, even of material tempered
above 450°C, it may be presumed that the prevention of corrosion allowed simple mechanical hydrogen-
induced fracture, and that the intergranular path taken without polarization or with anodic polarization is
due to an APC mechanism which may not only dictate the brittle-fracture path but also accelerate
hydrogen production. The active path component is not necessary to explain the effect of tempering in
all cases. The same pattern is obvious for the notched simple hydrochloric acid test (Fig. 7) as for the
neutral chloride solution tests (Fig. 5) and the selective corrosion mechanism is not applicable to acid
solutions but the relative differences between times for, say, 250°C tempered and 450°C tempered
samples in the two test media should be noted. With the extremely aggressive sulphide cracking test
medium, no "trough" due to tempering at intermediate temperatures is apparent (Fig. 8) and the effect of
tempering on the time rupture may be interpreted simply in terms of strength and toughness. Under some
circumstances, then, there may be an active path component in the mechanism, depending upon
treatment and environment, although an exclusive active path mechanism is unlikely.
Commercial 13% chromium steels usually contain between 11,5 and 14% chromium. There is
little evidence (Fig. 19) that variations over a much wider range (1 to 12%) have any significant effect on
Kiscc at least of lightly tempered steel. Effects of chromium content variation on the time to rupture of
precracked samples at a given stress intensity vary with environmental conditions (Fig. 20).
The influence of tempering temperatures on the SCC susceptibility of 13% Cr martensitic steels
with different contents of Ni and C was also studied by Ozaki and Ishikawa'24'13'. The testing
environment was high purity oxygenated water at temperatures in the range 150 to 288°C using slow
strain rate technique. The water was pH 6,5 , dissolved oxygen (DO) 8 ppm and the conductivity prior to
the test l|is/cm. The tensile strain rate of 10" /s was used. The SCC susceptibility is defined as the ratio
18
of the SCC fracture area Sscc /(Sscc+Sd). Intergranular corrosion depth was evaluated by immersion in
6,8% HNO3 solution at room temperature for 56 hours.
Figure 21 shows the summary of tensile strength and the intergranular corrosion depth as a
function of tempering temperature and Ni content. The tensile strength decreases with increasing
tempering temperature from 400°C to 600°C. For the steels containing 3,5 and 5wt% Ni the strength
increases again by tempering at 700°C. The impact strength and elongation increase monotonically with
temperature and Ni content. The IGC depth peaks in the tempering temperature range of 500 to 650°C,
showing the sensitization of the steels. The width of the tempering temperature range showing the
sensitization increases with increasing Ni content. Aci transformation point lowers with Ni addition and
this makes the selection of tempering temperature range for the steels with higher Ni content rather
limited, that is, below 600°C.
Figure 22 shows the SCC susceptibility in the high purity water at 200°C as a function of
tempering temperature and Ni content. The tempering below 500°C yields a high SCC susceptibility
regardless of Ni content. The SCC behavior becomes complex above 550°C.The steel with lower Ni
content shows inmunity to SCC when tempered at 600-700°C, while the steels with higher Ni content
show some susceptibility to SCC across whole tempering range.
The effect of hardness and IGC susceptibility on the SCC behavior of the steels examined in this
work is summarizes in figure 23. The IGSCC behavior is closely related to the IGC depth, while the HE
behavior is related to the hardness. The steel with low carbon and low Ni contents, tempered at higher
temperature so as to reduce the hardness to less than HV 280 and IGC depth to less than 20-50|im, is
immune to SCC. This steel can be concluded to be highly resistant to SCC judging from the severe stress
condition applied under slow strain rate test.
Figure 24 shows the guideline for the SCC-free steel design. This summarizes the SCC and HE
behaviors as a function of hardness and IGC depth for the steels with different C and Ni contents. The
figures in the diagram indicate the tempering temperature. The steel with high C and low Ni shows a
high IGC susceptibility in the tempering temperature range of 450°C to 700°C and no SCC-free zone can
be found. The steel with low C and low Ni shows a lower IGC susceptibility and lower hardness when
tempered above 650°C and the SCC-free zone can be readily defined, while the steel with low C and
high Ni shows a rather complex behavior, particulary above 600°C and the SCC-free tempering zone is
rather difficult to establish.
19
The SCC-free zone as a function of C and Ni contents and tempering temperature, determined at
288°C can be seen in figure 25. The SCC free zone can be clearly established for the steel tempered at
600-620°C and 650~660°C, but not for the one tempered at 55O~58O°C. The SCC-free zone can be
found when both the Ni and C contents are low. A Ni content of less than 4% is recommended for a C
content of less than 0,08% and Ni must be lower than 2,4% when C is less than 0,17%.
The effect of the tempering temperature is also shown in the SCC tests performed by I.L. Wilson
et al(~ ' on samples of martensitic, austenitic and duplex stainless steels. The martensitic steel samples
were AISI 410 in the form of C-rings. They were stressed to 240 MPa, to 90% of yield and with plastic
deformation. Several testing environments were used "reference boiler water chemistry", NaOH (19%
and 50%), PbO and Hg contaminants and Cl* (100 ppm) with sporadic oxygen additions, at 332°C. Part
of the results are summarized in table 9, below.
Sample Condition
As received
Tempered 650°C
Tempered 565°C
Tempered 482°C
Ref. Boiler Water
35000 h,
240 Mpa. Plast.def.
NCll)
NC
NC
NC
Ref. Boiler Water
35000 h,
0,15a.v
NC
NC
SCC
SCC
10% NaOH
4800 h, 0,9CTV
NC
NC
NC
SCC
Cl" (100 ppm)
10600h, 0,9ay
NC
NC
NC
SCC
(1)NC = No cracking
Table 9: Results of SCC tests of type 410 steel samples.
A correlation between SCC and chromium depletion at the prior austenite grain boundaries was
established in the work by P. Doig et a l u / . Samples were taken from a tube of a commercial 12% Cr Mo
V martensitic steel (11,5 Cr, 0,84 Mo, 0,51 Ni, 0,59 Mn, 0,22 V, 0,27 Si, 0,22 C). Austenitizing was
perfomed at 1100°C and tempering at various temperatures in the range 500 to 750°C. The samples were
stressed by bending and immersed for 106 s in a boiling deareated solution of 0,01 M NaOH plus 0,lM
NaCl. The conclusion, of this work was that the susceptibility of quenched and tempered 12% CrMoV
martensitic stainless steel to stress corrosion cracking in alkaline chloride solution is related to the
existance of a continuous chromium depleted concentration profile around the prior austenite grain
boundaries. Removal of this susceptibility occurs when this profile is destroyed by overlapping diffusion
fields from coarsening M23C6 precipitates in the prior austenite and martensitic interlath grain
boundaries. The tempering heat treatment necessary to eliminate susceptibility does not result in a
20
constant material hardness but rather a value which decreases as the tempering temperature increases in
the range 500 to 750°C. For the present steel, the criterion for removing susceptibility to stress corrosion
cracking, on the basis of tempering to a hardness value of < 280 HV10, is adequate to ensure immunity.
It is not correct to assume, however, that a similar criterion may be applied to other steels since it is not
directly based on that parameter which is responsible for the cracking susceptibility. Similar/, it is not
necessarily correct to conclude that hardness values > 280 HV10 will represent a susceptible tempering
condition.
pH=2
Eapp = 640
mVH
pH = 7
Eapp = 540
m\'„
pH =10
Eüpp = 540
mV,,
T° (°C)
25
50
75
100
25
50
75
100
25
50
75
100
tr(h)
57
56
26
17
59
56
24
20
60
53
34
22
UTS
(ksi)
105
107
101
89,4
109.5
104
97
103
110
108
95.6
103
CTy
(ksi)
68
71
81
81,3
59.7
67
83
96.5
65
76
74.4
66
Elongation
20,52
20.16
9.36
6.12
21,24
20.16
8.64
7.20
21.60
19.08
12.24
7.92
Type of Fracture
Ductile
cup & cone
Ductile
cup & cone
Brittle
IGSCC&TGSCC
Brittle
IGSCC&TGSCC
Ductile
cup & cone
Ductile
cup & cone
Brittle
IGSCC&TGSCC
Brittle
IGSCC&TGSCC
Ductile
cup & cone
Ductile
cup & cone
Brittle
IGSCC&TGSCC
Brittle
IGSCC&TGSCC
Surface
Appearence
Black-
General
Corrosion
Pitting
Pitting
Clean
Clean
Covered with
Film
Pitting
Clean
Covered with
Film
Pitting
Pitting
Aver,
Current
Density,
mA/cnr
+ 0.008
+0,12
+0,89
+2,47
0.0045
+0.009
+0,114
+1.86
+0,004
+0.013
+0.86
+0.543
Table 10: Stress Corrosion Test Result for type 403 Stainless Steel in 0,0JM Na2SO4, pH 2.7 and 10-6 -1at Various Temperatures. Starin rate = 10" s" .
21
Stress corrosion cracking of samples of AISI type 403 martensitic steel in a 0,01 M N
solution was observed by Bavarian et al(:8). The samples were austenitized at 960°C and tempered at
650°C. Tests were perfomed by means of SSRT in solutions with pH= 2, 7 and 10. The results of the
stress corrosion tests are collected in the table 10. These data show that intergranular SCC occurred at
temperatures of 75 and 100°C, but not at 25 and 50°C. The time to failure is approximately the same for
all solutions, regardless of the pH. Significant sulfur contamination of the surface oxide films has been
observed using Auger Electron Spectroscopy depth profiling. This does not parallel the trend in the
susceptibility of the film to localized breakdown. Accordingly, it is apparent that sulfur contaminations
from the solution is not the prime cause of localized attack on this steel. Nonmetallic inclusions of MnS
and chromium carbides are the most susceptible sites for pit nucleation. The corrosion attack usually
starts at the boundaries between inclusions and the passivated metal. This leads to localized attack in the
forms of pits, which then act as sites for crack nucleation.
(29)An effect of heat treatments on samples of martensitic steels has been reported by Tsubota et al
The samples were prepared from martensitic steels of the following designations: CA6NM (13% Cr, 4%
Ni), SUS type 431 (17Cr, 2,4Ni, 0,2C) and type 630, although these are not of direct interest to fusion
application. The chemical composition of the alloys are listed in Table 11. The CBB (Crevice Bent
Beam) was employed for the SCC tests. The tests were carried out at 288°C for 500 hours, and the SCC
susceptibility was evaluated by crack depth measurement on longitudinal section of the specimen.
Alloy
403
420J1
CA40
F6NM
431
630
C
0,14
0,16
0,26
0,033
0,18
0,04
Si
0,44
0,29
0,49
0,36
0,82
0,29
Mn
0,74
0,39
0,77
0,65
0,80
0,74
P
0,028
0,031
0,02
0,02
0,036
0,027
S
0,007
0,013
0,01
0,013
0,011
0,003
Ni
0,25
0,30
0,10
3,87
2,42
4,06
Cr
11,70
12,12
11,9
12,74
16,98
15,71
Others
-
-
-
Mo 0,50
-
Cu 3,24
Nb 0,28
Table 11: Chemical composition of the alloys investigated (wt%)
22
The alloys examined, except for CA40, were prepared as forged bars and heat treated with
following conditions:
403: 1050°C/3h + 500, 550, 600, 650, 700, 750, 800°C / 8h.
420J1: 950°C/4h + 500, 550, 600, 650, 700, 750, 800°C / 8h.
CA40: 950°C/4h + 500, 550, 600, 650, 700, 750, 800°C / 8h.
F6NM: 1100°C /5h + 450, 500, 530, 550, 580, 600, 650°C / 8h.
431: 1050°C/5h + 500, 550, 600, 630, 650, 700, 750°C/ 8h.
630: 1038°C /0,5 + 480, 550, 565, 580, 600, 620°C / 5h.
The average of the maximum stress corrosion crack depths observed in ten specimens of each heat
treated alloy are shown in Fig. 26. As-quenched and low temperature tempered specimens showed high
susceptibility. The SCC susceptibility of martensitic stainless steels is related to their hardness and
tensile strength. Martensitic stainless steels with HV 340 or au > 110 Kg/mm" posses high
susceptibilities in a high temperature environment, regardless of the alloy specifications. Martensitic
stainless steels should be well-tempered and anion concentration in the water must be kept as low as
possible.
The influence of hardness levels in the susceptibility to SCC, has also been recently studied
by Lapeña et al a }. In this work the material tested was the low activation martensitic steel F-82H
modified, considered as possible structural material for fusion applications. The nominal composition
of this alloy is 7.65Cr, 2.1W, 0.100C, 0.16Mn, 0.14V, 0.003S, 0.002Ta wt%, balance Fe. Crack
growth rate tests were carried out in two different material states (see table 12), using compact tension
specimens (CT) 12 mm thickness. Previous to crack growth corrosion tests, all specimens were
precracked in air at frequency of 22 Hz and R = 0.1. The samples were tested under constant load.
The water temperature was 260°C, with 0,27 ppm of lithium as additive and 2 ppm of hydrogen at
room temperature.
Material condition
Normalized at 1075°C/30'
Normalized at 1040°C/30' + Tempered at 750°C /I h air cooled.
HV30
405
204
Table 12: Material condition and hardness of F-82H modified.
The results of this work showed important differences on crack growth rates in function of
hardness levels. The specimens normalized plus tempered (1040°C/30' + 750°C /I h air cooled) with a
23
hardness values of 204 not showed any crack growth after 1200 and 3225 hours testing. In contrast of
these results, the material tested only in the normalized state (1075°C/30') showed large crack growth
even in very short periods (= 1 mm during 4 hours). Other important results obtained by these
investigators was that some of the specimens broke during the tests (in some cases <12 hours),
indicating a fracture toughness values between 120 and 133 MPaVm. All the samples tested in this
material condition exhibited intergranular fracture. In this work, tests were also performed without
hydrogen in order to determinate the influence of hydrogen in the behaviour of this steel to SCC. The
crack growth rates obtained in both cases (with and without hydrogen) were in the same order the
magnitude as can be seen in figure 27. The F-82H modified with a hardness level of 405 showedQ
IGSCC and high crack growth rates of approximately 7.10" in the range of stress intensity factor
between 40 and 80 MPaVm. However, the steel in the normalized plus tempered state was not
susceptible to SCC.
The observed behaviour of F-82H modified to SCC is in agreement with the results obtained
by Tsubota<29> in which work, as mentioned previously, the samples tested on as-quenched condition
(high hardness level) or inappropriately tempered presented high susceptibility to stress corrosion
cracking. Both studies seem to indicate that in the case of martensitic steels a closely relation exist
between hardness level and susceptibility to SCC.
Tsubota also mentioned the possible influence of hydrogen produced during the corrosion
process. He measured the amount of hydrogen in the steel and find between 3.5 and 9 ppm (Fig.28),
but a relationship between crack depths and hydrogen, as can be observed between crack depths vs.
hardness or crack depths vs. tempering temperature (Figs. 26and 28), can not be clearly seen.
Lapeña(1 ', based in the studies of Boler et al( ' , suppose about 0.004 ppm hydrogen in the material.
All these authors mentioned that the SCC behaviour is very similar to hydrogen embrittlement
behaviour.
Although is not usual to use the martensitic steel in as-quenched or normalized conditions, the
hardness levels obtained by these heat treatments can be representative of heat affected zone (HAZ)
hardening after unsuitable post-weld heat treatment.
The results of these works are very important also for the low activation martensitic steels
since the point of view to optimize the temper temperature in order to obtain a good jointly
mechanical and corrosion properties.
24
5.- IRRADIATION-ASSISTED STRESS CORROSION CRACKING (IASCC)
Irradiation assisted stress corrosion cracking has been used to describe intergranular
environmental cracking of materials exposed to ionizing irradiation/"". While more restrictive
interpretations have been applied, a consensus has developed that the term IASCC be applied to all
instances where environmental cracking has been accelerated by radiation, whether it acts singly or
jointly to alter water chemistry, material microchemistry, material hardness, creep behavior etc.
The effects of radiation on material properties have been widely recognized and studied for
decades, although the early emphasis was on radiation hardening, swelling and creep(J" }, and water
chemistry1'3 ' '. While the possible effects of radiation on environmentaly assisted cracking are
numerous, many are poorly quantified and /or their effect on IASCC is completely unknown. The list of
radiation phenomena which are potentially important includes: i) radiation induced segregation; ii)
radiation elevation of crack tip and crack mouth corrosion potential; iii) radical and ionic species (e.g.,
H2 O2 , OH, HO2, e"aq); iv) transmutation to form species which are, e.g., soluble in the crack solution
(e.g., NO3 ) or embrittling to the material (e.g., H); v) Radiation-enhanced creep-relaxation; vi)
radiation hardening; vii) microscopic and macroscopic swelling; etc. From such a large list of complex
phenomena, it is necessary, as a first step, to identify those factors which are likely to have the primary
impact on environmental cracking susceptibility. For this reason, and since insufficient data are available
on direct effects on cracking of, e.g., radiation induced creep and radiation hardening, it has been
decided by international consensus that IASCC should be identified and studied taking into account: a)
micro-compositional changes from radiation induced segregation, and b) corrosion potential elevation
from oxidizing species produced by gamma and neutron interaction with water.
No reports of IASCC failures of martensitic stainless steels appear to be published in the open
literature. Post-irradiation SCC tests of specimens from steel of type 1.4914/MANET, 1.4914 with a
laser weld, FV448 and reduced activation alloy LA12TaLC (Table 3) were perfomed by Nystrand ( ' J '.
Corrosion coupons of the steels had been previously tested for 1460 or 4947 hours respectively (time at
a reactor power of > 30 MW) in water at a temperature of 275± 10°C and a pressure of 90 bar. The
displacement doses were estimated to be 0,60 and 2,34 dpa for the respective exposure times. The SCC
tests were perfomed in deionized water satured with air at room temperature and filtered to remove
humic matter. The main part of the tests were perfomed under 3-point bending at an initial stress of 95%
of yield, in a low flow autoclave. Two of the irradiated specimens were provided with a transverse notch
at the position of maximum strain and were stressed in the specimen holder until a small permanent
deformation occurred. The total testing time for each set of specimens was 1500 hrs. Visual and
metallographic examinations failed to reveal any cracks in the specimens tested. The structure of all the
25
specimens was tempered martensite.
As mentioned in this section, one of the important subject related with IASCC is the influence of the
microcompositional changes at grain boundaries under irradiation. For this reason, we have believed
necessary to intoduce in this report one section dedicate to know the segregation behaviour of
ferritic/martensitics steels irradiated.
6.- RADIATION-INDUCED SEGREGATION (RIS)
Several investigations of radiation-induced segregation in martensitic and ferritic steels have been
perfomed with the main purpose to provide information about the effects on the mechanical properties.
T.S. Morgan et al<38) examined specimens of annealed and tempered FV448 after irradiation to
high neutron doses. One set of specimens was irradiated to 25 dpa at 400°C, another to 46 dpa at 465°C.
Concentration profiles at lath boundaries were determined by means of STEM. The results, similar for
both irradiation conditions and showed prominent enrichment of Ni and depletion of Fe at a lath
boundary. Cr showed a w-type concentration profile.
R.E Clausing et al<39) sudied the segregation effects in neutron irradiated type HT-9 martensitic
steel. Annealed (1035°C for lh / air cooled) and tempered (760°C for 1 h) specimens were irradiated to
about 13 dpa at 410, 520 and 565°C. Control specimens were aged at the appropiate temperatures for
periods of 15000 h to match the times for the irradiated specimens. Following irradiation or thermal
aging, each specimen was inserted in a special fracture device Auger spectrometer and its was cooled
approximately at -196°C and fractured by impact. All specimens broke well below the Ductile-Brittle-
Transition-Temperature (DBTT) and exhibited a predominantly brittle fracture. Randomly oriented
micro-facets of the order 2-5 um with evidence of ductile-tearing at facet edges were observed in the
SEM Stereomicrographs. These micro-facets are believed to be related to the underlying lath packet
structure. In addition to the overall micro-facetted structure, several much larger facets of the order 25-50
urn diameter were present on the surface. These macro-facets were very smooth and clearly represented
crack propagation along a different microstructural component. In specimens irradiated at 520°C and
565°C very few of these macro-facets could be found, and in the thermal controls none could be found.
However, a few such regions were observed in thermal controls, which had been previously first
hydrogen charged and then broken at room temperature. .
26
For all the specimens, AES analysis of the micro-facetted regions yielded elemental
concentrations at the same levels as the bulk chemical analyses, indicating that there had been no
segregation in these regions. The macro-facet on the hydrogen charged thermal controls showed
enrichment of C. Cr and Mo, suggesting that the fracture occurred through a carbide rich region. The
macro-facets in the specimen irradiated at 410°C displayed significant levels of segregation of Ni, Cr, Si
and P (Fig 29). At the fracture surface, a chromium enrichment of a factor of ~ 1,4 above its bulk
concentration and silicon enrichment by a factor of ~ 10 were observed. The concentration of
phosphorus was ~ 1 at. % representing an enrichment over the average concentration by a factor of ~
100. In each case, these concentrations decreased with depth approaching bulk values within ~ 30 nm
below the surface. Nickel showed a completely different behavior. The measured surface concentration
of nickel was ~ 4 at.%, i.e., about 8 times higher than the average bulk concentration. The concentration
of nickel increased with increasing depth, the concentration of nickel increased, reaching a value of ~8 at
% at 60 nm. These authors concluded that the radiation induced segregation of Cr, Ni, Si, and P at
410°C, but not in specimens irradiated at higher tempeartures or in thermal controls. The concentrations
of Cr, Si and P decrease rapidly with increasing distance form the boundary. However, in contrast, the
concentration of Ni increases initially and then persists at high concentartion over a distance of a least
120 nm. In this work, the authors do not describe the causes of the behavior of the Ni.
Compositional changes have been measured at grain boundaries, dislocation loop, and precipitates
in Japanese Ferritic/Martensitic Steels (JFMS) (Fe-S^Cr^Mo-O^Ni-OJSi-O^Mn-OJNb-OJV-O.OSC)
and its three high purity model alloys, namely Fe-10Cr, Fe-10Cr-lNi and Fe-10Cr-5Ni by T. Muroga et
al( '. These specimens were normalized at 1050°C for 1 h and tempered at 750°C for 2 h. Irradiations
and in-situ microstructural observations were carried out with 1 or 1.25MeV electrons in HVEM. The
investigations showed an enrichment of silicon and depletion of chromium at grain boundaries (Fig. 30)
and precipitate-matrix interfaces in Japanese Ferritic/Martensitic Steels irradiated at 500°C. Similar
depletion of chromium was observed at grain boundaries in Fe-10Cr and Fe-10Cr-lNi alloys irradiated
above 300°C. Also the radiation induced decomposition of precipitate-matrix interfaces in JFMS has
been measured. In this steel, M23C6, M&C, and MC type precipitates, identified by EDS and
microdifraction, are observed. An example of the solute concentration profile before and after irradiation
for a MûC type precipitate is shown in Fig. 31. The change in solute concentration takes place mostly
within 50 nm from the precipitate-matrix interfaces. In order to show clearly the manner of segregation,
the change of concentration in matrix 25 nm from the interface and the precipitate 50 nm from the
interface are indicated in Figs. 32 and 33 respectively. In these figures, the result for molybdenum is not
included, as the change in concentration by irradiation is small. Figure 32 shows that, for any precipitate,
silicon is enriched and chromium and nickel are depleted at the matrix near the interface. This
27
decomposition is qualitatively the same as that at grain boundaries shown in figure 30. However, the
depletion of nickel is more prominent than that at grain boundaries, especially for M23C6, MC
precipitates. Figure 33 indicates that the change in silicon and chromium concentrations in the
precipitates takes place in a similar manner to that observed at nearby matrices. These experiments, the
chromium depletion and silicon enrichment are obtained at defect permanent sinks, namely, grain
boundaries and precipitate-matrix interfaces. These results apparently obey the conventional solute size
dependence. On the contrary, the behavior of nickel is rather complex. The characteristic nickel
segregation may be expalined by the interaction of nickel with both intersticials and vacancies. Namely,
the association of nickel with insterstitials can result in the enrichment of nickel at interstitial-biased
sinks and nickel transportation in the opposite direction to the vacancy flow can cancel the
decomposition at neutral sinks and change reversely the decomposition at vacancy-biased sinks.
Depletion of Cr at grain boundaries was also observed by H. Takashashi et al( } in alloys of Fe-
5Cr and Fe-13Cr after irradiation with 650 KeV electrons to 3 dpa at 400°C. When these material were
irradiated, no radiation induced precipitates or voids were nucleated. Therefore, the compositional
analysis was perfomed only in the irradiated region including a grain boundary. Figure 34 shows an
example of the results obtained after irradiation at 400°C to a dose of 3 dpa. The concentration of Cr in
Fe decreased at or near the grain boundary region within 200 nm from the grain boundary than at the
matrix. The concentration gradient of Cr near the grain boundary became steeper with increasing
concentration. Also, a Cr enriched zone was formed around the depleted zone. The results obtained by
Takashashi are in contrast with S. Ohnuki(42), who found strong enrichment of Cr at the grain boundaries
after irradiation of specimens of Fe-13Cr and Fe-13Cr-lSi alloys with 200 KeV carbon ions to a dose of
57 dpa at 525°C. A third alloy Fe-13Cr-lTi showed uniform Cr concentration across the grain
boundaries after irradiation.
The enrichment of Si at grain boundaries also was observed by Kimura et al (4j'44) . In this work
grain boundary chemistries in low activation 9%Cr- 2%Mn-l%W and 12%Cr-6%Mn-l%W steels were
measured by AES (Auger electron spectroscopy) after irradiation in the FFTF/MOTA at 365°C up to
doses of 10 and 25 dpa. In the 9%Cr alloy, grain boundary segregation of Si was deteced, but no
significant changes of the other elements were observed. However, in the 12%Cr steel, Si and Mn
segregation was reconogized after irradiation to 25 dpa. In both materials, little significant effect of
irradiation on P and S segragtion was observed. Figure 35 show the dependence of grain boundary
concentration of Mn and Si on the irradiation dose for both steels. As can be seen in the graph, Si
segregation in the 12% Cr alloy significantly increased with increasing the dose from 10 dpa to 25 dpa,
while that the Mn appears to decrease. In contrast to this, grain boundary concentartion of Si in the 9%
28
Cr alloy does not change by the increase in irradiation dose. Both steels showed intergranular fracture
after irradiation at 365°C to 10 and 25 dpa.
According to Muroga ( ' and Takashashi( ' , recent studies performed by Schäublin et al ' J ) also show
Cr depletion at the grain boundaries. Schäublin {3) investigated the segregration behaviour to irradiation
of the low activation F-82H modified ferritic /martensitic steel using energy filtered transmission
electron microscopy (EFTEM). The material was irradiated with 590MeV protons in the PBREX facility
to a dose of 0,5 dpa at 250°C. With this technique and after the irradiation, the F-82H modified showed
Cr and Fe depletion at the grain boundaries (prior austenite grains and lath martensite boundaries).
29
7.- SUMMARY AND CONCLUSIONS
General corrosion of high chromium ferritic/martensitic steels in high temperature water (300°C-
350°C) results in the formation of double oxide layers. The outer layer consists of magnetite crystals and
the inner layer consists, in general, of a fine-grained chromium iron spinel structure which controls the
rate of corrosion. These types of steels under water chemistry conditions of LMFBR and PWR show
moderate corrosion rates and parabolic corrosion rate law at temperatures up to 350°C.
The effect of irradiation on corrosion has been studied for different martensitic steels, candidates
for fusion applications, such as 1.4914 (MANET), FV 448, and reduced activation steels (LA7Ta,
LA12Ta and LA12TaLC). The results of this investigation showed that the corrosion of these steels at
275°C-300°C depends on the situation of the coupons in the reactor and the effects of water flow on the
corrosion kinetics. This may be the reason for the diferences found in the results reported on the one
hand by Källstrom(21>, Gott(22> and the other by Ashomore(1:>>.
The influence of tempering temperature on the SCC susceptibility of 13%Cr martensitic steels has
been studied with différents techniques (SSRT, CBB, etc) to evaluate the susceptibility of these alloys.
The results of the investigations realized by different investigators revealed that for these materials,
tempering in the range 350°C to =600°C is particular/ detrimental to the SCC resistance. Other authors
concluded that the criterion for removing susceptibility to stress corrosion cracking, on the basis of
tempering to a hardness value of <280HV10, is adequate to ensure immunity, althought, it is not
necesasrily correct to conclude that hardness values > 280HV10 will represent a susceptible tempering
condition.
No reports of IASCC failures of martensitic stainless steels appear to be published in the open
literature. Post irradiation SCC test have been perfomed on steel of type MANET, FV 448 and reduced
activation alloy LA12TaLC. In this work, there is no indication of irradiation assisted stress corrosion
cracking in the specimens after irradiation to 2,3 dpa. This aparent resistance to IASCC may be a
consequence of the fact that the steels were not sensitized by the irradiation.
30
Results of studies of radiation-induced segregation (RIS) effects in martensitic steels are of
interest, since one of the investigations of stress corrosion cracking'" ' showed a correlation between
SCC and depletion of Cr. The studies of RIS in martensitic and ferritic steels have, in general, not
indicated any depletion of Cr at prior austenite or ferrite grain boundaries. Depletion of Cr was observed
only in two cases, where the samples were irradiated with electrons. Martensitic steels are expected to be
resistant against radiation-induced segregation owing to the high density of point defect sinks in the
martensitic structure.
31
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Materials 191-194 (1992) p. 885-889.
45. Schäublin.R. Spätig.P, Victoria.M. "Chemical segregation behaviour of the low activation
ferritic/martensitic steel F-82H". Journal of Nuclear Materials 258-263(1998) p. 1350-1355.
36
LIST OF FIGURES
Figure 2: Metall loss versus time for 21/4 Cr and 9% Cr ferritic steels. Average line is least squares fit
to data. (From Ref. 7).
Figure 3: Effect of heat flux and deposition of a-FezO:, on metal loss. (From Ref. 7).
(a) AVT water chemistry; low dissolved O2 (<10 ppb); 4100h.
(b) AVT water chemistry; high dissolved O2 (56 ppb); 3500 h.
Figure 4: Weight loss versus time for base and weld material of F-82H modified. (From Ref. 16-17)
Figure 5: Time to rupture versus tempering-temperature plots for three 13% Cr steels, smooth
specimens tested in 3%C1 solution at ambient temperature; initial stresses 695 or
1004MNni2. (From Ref. 23).
Figure 6: Time to rupture versus tempering temperature plots for smooths specimens of 13% Cr
steels tested in boiling 3% NaCl solution; initial stress 464 MNm"2. (From Ref. 23).
Figure 7: Effect of tempering temperature on time to rupture of notched samples of three 13% Cr
steels stressed in 1% HC1 at ambient temperature. (From Ref. 23).
Figure 8: Effect of tempering temperature on time to rupture of 13% Cr steels stressed in a sulphide
environment. Smooth samples stressed initially at 464 Mnm~~. (From Ref. 23).
Figure 9: Efects of tempering on hardness, impact strength, and corrosion rate in 10% HNO3 at 20°C
of hardened 0,28-13Cr steel. (From Ref. 23).
Figure 10: Effect of strenght (UTS) on time to rupture of 13% Cr steels of variying carbon content
tempered at 250°C, 450 °C and 550°C; smooth specimens were tested in 3% NaCl solution
at ambient temperature. (From Ref. 23).
Figure 11: Effect of strength (UTS) on time to rupture of 13% Cr steels of varying carbon content
tempered at 250°, 450°, or 550°C; smooth specimens were tested exposed to an industrial
atmosphere. (From Ref. 23).
37
Figure 12: Effects of galvanostatic polarization on time to rupture of a 13% Cr steel tempered at
250°C stressed in 3% NaCl solution at ambient temperature. (From Ref. 23).
Figure 13: Effect of initial applied stress on time to rupture of 13% Cr steels hardened and tempered at
250°C; corrodent used was 3% NaCl solution at ambient temperature. (From Ref. 23).
Figure 14: Effect of initial applied stress (Smooth specimens) on time to rupture of 13%Cr steels
hardened and tempered at 250°C, corrodent used was an industrial atmosphere. (From Ref.
23).
Figure 15: Initial stress intensity versus time to rupture plots for three 13%Cr steels hardened and
tempered at 250°C; precracked samples tested in 3%NaCI solution at ambient temperature.
(From Ref. 23).
Figure 16: Effect of polarization on crack growth rate of 0,21C-13%Cr steel hardened and tempered at
250°C; precracked specimens were tested in 3% NaCl solution at ambient temperature
unpolarized, polarized anodically, or polarized cathodically. (From Ref. 23).
Figure 17: Initial stress intensity versus time to rupture plots for O,21C-13Cr steel hardened and
tempered at 250°C; precracked specimens were tested in NaCl solutions of various
concentrations at ambient temperature. (From Ref. 23).
Figure 18: Influence of applied potential (Potenciostatic control) on time to rupture of precracked
specimens of 0,21C-13Cr steel, hardened and tempered at 250°C, in 3%NaCl solution at
ambient temperature. (From Ref. 23).
Figure 19: Initial stress intensity versus time to rupture curves for ~ 0,3% carbon steels of 0-12% Cr
content tested in 1050°C and 200°C-250°C tempered condition; precracked samples were
tested in 3% NaCl solution at ambient temperature. (From Ref. 23).
Figure 20: Influence of polarization (galvanostatic) on times to rupture of -0,3% carbon steels of
12%Cr content tested at initial stress intensity of 3362 KgcrrT " in NaCl solution at ambient
temperature; precracked samples not polarized and polarized anodically or cathodically at
15mAcm"2; steels hardened and tempered at 200°-250°C. (From Ref. 23).
38
Figure 21: Summary of tensile strenght and IGC depth for the 13% Cr steels with typical Ni content as
a function tempering temperature. (From Ref. 24).
Figure 22: Dependence of the SCC susceptibility in the high purity water at 150°C on tempering
conditions for the 13% Cr steels with typical C and Ni content. (From Ref. 24).
Figure 23: The critical SCC susceptibility-hardness-IGC depth diagram. (From Ref. 24).
Figure 24: The critical SCC susceptibility-hardness-IGC depth diagram for the typical C and Ni
contents. (From Ref. 24).
Figure 25: The critical SCC susceptibility C and Ni content diagram for the typical tempering
temperatures. (From Ref. 24).
Figure 26: Effect on tempering temperatures on SCC crack depths of alloys investigated. SCC test
was conducted in high temperature (288°C) water by CBB technique. (From Ref. 29).
Figure 27: Crack growth rate versus stress intensity factor of F-82H modified on normalized
(1075°C/30') state. (From Ref. 16).
Figure 28: Hydrogen concentration measured in specimens after CBB test in 288°C water for 500
hours. (From Ref. 29).
Figure 29: Concentration versus depth from the surface fracture macro-facets on the specimen
irradiated at 410°C. The concentrations are the average values from seven areas on three
facets. (From Ref. 39).
Figure 30: The solute concentration at and near a grain boundary in irradiated Japanese Ferritic
/Martensitic steels. The data points are the average of five to eight measurements at the
same distance from the grain boundary. The error bars connect the minimum and maximum
values for the five to eight cases. (From Ref. 40).
Figure 31: The compositional profile near a MoC type precipitate-matrix interface in Japanese
Ferritic/Martensitic steels before and after irradiation at 500°C. (From Ref. 40).
39
Figure 32: The change in solute concentration at matrix in the vicinity of the precipitates in Japanese
Ferritic/Martensitic Steels by irradiation at 500°C. (From Ref. 40).
Figure 33: Same as Figure 32 except at the precipitates in the vicinity of precipitate-matrix interface.
(From Ref. 40).
Figure 34: Concentration profile of Cr in Fe at near a grain boundary as a function of distance from the
grain boundary. (From Ref. 40).
Figura 35: Dependences of grain boundary concentration of Si and Mn on the irradiation dose in (a)
9%Cr-2%Mn-l%W and (b) 12%Cr-6%Mn-l%W alloys. (From Ref. 43).
40
leu {»ni
Figure 2: Metal loss versus time for 21/4 Cr and 9%
Cr ferritic steels. Average line is least squares fit to
data. (From Ref.7).
«n ia i ai4X)SSj3 NOM «Jira uf í«l
rto*3a *o s
? •
i Jim»- • un* ivi*t~
- 1 i
g
E
SS
OI ]
cO)
90
80
70
60
50
40
30
a Base metal (TIG)o Base metal (EB)c Weldment metal (TIG)O Weldment metal (EB)
1000 2000 3000 4000
Time (hours)
5000
Figura 4: Weight loss versus time for base and weld
material of F-82H modified. (From Ref. 16 and 17)
1 0 a
10'
!¿Í1O'
1 O -
200 300 •a OO 5O0 6OO 700
1CXX MN m"1 69S MN m"2
TEMPERING TEMPERATURE, *C
Figure 3: Effect of heat flux and deposition of a-
Fe;O3 on metal loss. (From Ref.7).
(a) AVT water chemistry; low dissolved O2 (<10
ppb);4100h.
(b) AVT water chemistry; high dissolved O2 (56
ppb); 3500 h.
Figure 5: Time to rupture versus tempering-
temperature plots for three 13% Cr steels; smooth
specimens tested in 37c NaCl solution at ambient
temperature; initial stresses 695 or 1004 MNm"2.
(From Ref.23).
41
IC
O
,10'
x O O7-/.Co O 26-/.CD O 32-/.C
2OO 3OO "JOO 5OO 6OO 7OOTEMPERING T E M P E R A T U R E , *C
2OC 3OO -aOOTEMPERING
Figure 6: Time to rupture versus tempering
temperature plots for smooths specimens of 13% Cr
steels tested in boiling 3% NaCl solution; initial
stress 464 MNni2. (From Ref. 23).
Figure 8: Effect of tempering temperature on time to
rupture of 13% Cr steels stressed in a sulphide
environment. Smooth samples stressed initially at 464
Mnm":. (From Ref. 23).
TO-
10"'
O19--C I AC 1O5O-COZB-Í.CI
2OC 3OQ &DC 7OC
10O¿ MNrrr' 695 MNn-;*TEMPERING TEMPERATURE (In.1, "C
200
—1
TH
,
1»LUCy-
*J4h—
U<c.27û0
600<
0n=500
ci
2^00Û
tr$""300
0Vù
naronessIiodcorrosion rate
-3-Ou,-
0 100 200 300 400 500 600700TEMPERATURE, *C
Figure 7: Effect of tempering temperature on time to
rupture of notched samples of three 13% Cr steels
stressed in 1% HC1 at ambient temperature. (From
Ref. 23).
Figure 9: Effects of tempering on hardness, impact
strength, and corrosion rate in 10% HNO3 at 20=C of
hardened 0.28-13Cr steel. (From Ref. 23).
42
1700-
tempérée,
x 250o -áSOù 550
I Tstress, MNrrf1
TOO!10O4695
TIME TO RUPTURE, h
10'
no
105O"/250-C condîfon3*/. Na Cl
cast Oplain tensile speamensat 1OOOMN m'2
O
10
10 _L
.adCnarpynens at 22MNm 2
_ I . I-16 - 1 2 - 6 - 4 0 4 8 12
4 vo -ve e-APPLIED CURRENT DENSITY, mA cm"2
Figure 10: Effect of strenght (UTS) on time to
rupture of 13% Cr steels of variying carbon content
tempered at 250°C, 450 °C and 55OQC; smooth
specimens were tested in 3% NaCl solution at
ambient temperature. (From Ref. 23).
Figure 12: Effects of galvanostatic polarization on
time to rupture of a 13% Cr steel tempered at 250°C
stressed in 3% NaCl solution at ambient temperature.
(From Ref. 23).
o \ tempered,°C stress,MNrr, X*X . OCA *r\r\j
1400
)0¿ 103 104 105
TIME TO RUPTURE , h
Figure 11: Effect of strength (UTS) on time to
rupture of 13% Cr steels of varying carbon content
tempered at 250°, 450°. or 550°C; smooth specimens
were tested exposed to an industrial atmosphere.
(From Ref. 23).
- 200h
x 0- ITÍ .C
o Q22*IJZD
ß :
AC, 1050"C-2nst250'C
10 102 103 104 105
TIME TO RUPTURE, n
Figure 13: Effect of initial aplied stress on time to
rupture of 13% Cr steels hardened and tempered at
250°C; corrodent used was 3% NaCl solution at
ambient temperature. (From Reñ 23)
43
1400'E2
x 01V/.C ]o 0 lev.C I AC 1050'C • 2h at 250-C
- D 028-/.C
V
UJ
< 600-2
2 0 0 -
10 103 10'TIME TO RUPTURE
10=
UJ
1 1 O " J
h-
OerO
V<
u
10"
1050*/250o conditionapplied
xDo
2000 4 000 6000 800010000INITIAL STRESS INTENSITY, kg cm"3'J
Figure 14: Effect of initial applied stress (Smooth
specimens) on time to rupture of 13%Cr steels
hardened and tempered at 250°C, corrodent used
was an industrial atmosphere. (From Ref. 23).
Eu01
I/)
I 1 I1050*/250* conditionoocn symbols 'valid' tests(with respect to specimen size)dosed symbols 'invalid' tests(with respect to specimen size)o 0T/.CD 0-27.CV 0-3°/.C
10" 1 10 Kr 1OJ
TIME TO RUPTURE, h
Figure 15: Initial stress intensity versus time to
rupture plots for thrree 13<7rCr steels hardened and
tempered at 250=C; precracked samples tested in
37rNaCl solution at ambient temperature. (From Ref.
23).
Figure 16: Effect of polarization on crack growth
rate of 0,21C-13%Cr steel hardened and tempered at
250°C; precracked specimens were tested in 3%
NaCl solution at ambient temperature unpolarized,
polarized anodically, or polarized cathodically.
(From Ref. 23).
5000
1050°/250°. condition
NaO,% pHo 03 7à 3-0 7x 100 7
Z Z
1 10 10' 1OJ
TIME TO RUPTURE, h
Figure 17: Initial stress intensity versus time to
rupture plots for 0,21C-13Cr steel hardened and
tempered at 250°C; precracked specimens were
tested in NaCl solutions of various concentrations at
room temperature. (From Ref. 23).
44
102
10
1
m-1
1O5O*/25O'C condition3*/. NaClprecracked Cnarpy specimensat 2239-6 kg cm"3#
/
\1 \ J\ 0 "*
\ -
-1200 -800 -400 0APPLIED POTENTIAL,mV(sce)
10' 1 10 102 1CTIME TO RUPTURE, h
TO"
Figure 18: Influence of applied potential
(Potenciostatic control) on time to rupture of
precracked specimens of 0,21C-13Cr steel, hardened
and tempered at 250°C, in 3%NaCl solution at
ambient temperature. (From Ref. 23).
Figure 19: Initial stress intensity -versus time to
rupture curves for ~ 0,39c carbon steels of 0-12% Cr
content tested in 1050°C and 200°C-250°C tempered
condition; precracked samples were tested in 3%
NaCl solution at ambient temperature. (From Ref. 23)
45
10-
-10'
10"
n i i r i i r
x unpoiari2edù cathodicaiiy polarized].,. -20 anodicaliy polarized J 1 D m A c m
0 1 2 3 4 5 6 7 8 9 10 11 12Cr, •/.
l$c « . . i « ¿ SO v m/56h
400 iCC tOO M • 500 ¡M ?0C 50C 60S !0C
Tempering !emp./"C x 5h
Figure 20: Influence of polarization (galvanostatic)
on times to rupture of ~0,37c carbon steels of 0-
12%Cr content tested at initial stress intensity of
3362 Kgcm"J~ in NaCl solution at ambient
temperature; precracked samples not polarized and
polarized anodically or cathodicaiiy at 15mAcni:;
steels hardened and tempered at 2O0°-250°C. (From
Ref. 23).
Figure 21: Sumamry of tensile strenght and IGC
depth for the 13% Cr steels with typical Ni content as
a function tempering temperature. (From Ref. 24).
46
too 5O0 60C
Temperinc temp./^C x 5h
7DC
Figure 22: Dependence of the SCC susceptibility in the high purity water at 150°C on tempering conditions for the
13% Cr steels with typical C and Ni content. (From Ref. 24).
X
E.oX!
lOÖCf
500
100
50
10
L --
-
_ i--
- t x .
•
N o
-
Oto O
. . . - s t . .
C
oooo
0
•
o ©
O ;
' o l eO Ä '
o
m
o
V-NV
O O CO
o!
- V "
I
e
j 2&8"C our« ««
j e :sccj D • No SCC
O
-^> HC #
¡1 o
200 25C 300
Haroness / M-
Figure 23: The critical SCC susceptibility-hardness-IGC depth diagram. (From Ref. 24).
47
\XK
B
X
cCJ
O
10* - M I 0.06C-5~6NÍ-13Cr I .
200 250 300 350 ¿00 ¿50 200
Hardness / Hv
250 300 350 ¿0C
Figure 24: The critical SCC susceptibility-hardness-IGC depth diagram for the typical C and Ni contents.
(From Ref. 24)
48
oo
(288 °C)
or. 4
2
6
2 r-
0
©
O
© o
SCC
'8
No SCC
• SCC
O no SCC
o fc
550~580°CTemoer
620*0Temper
650660Temper
0.05 0.1 0.15
C content /%
Figure 25: The critical SCC susceptibility C and Ni content diagram for the typical tempering temperatures.
(From Ref. 24).
49
£
"c.cOje
o
SC
C
2000
1500
1000
500
n
¡ 0—m
- A
X
co og
X
D0•
A
403
42CJ)
CA40
431
63C
As 0 500 60C
T e m p e r i n g T e m p e r a t u r e ( C )
700 800
Figure 26: Effect on tempering temperatures on SCC crack depths of alloys investigated. SCC test was
conducted in high temperature (288°C) water by CBB technique. (From Ref. 29).
50
1E-6
ë.s13
1E-7
Oen
o2ü
1E-8-
® with hydrogena without hydrogen
201 i '40
• i '60
' ' ' ' ' ' ' i ' ' ' ' ' ' '80
K (MPa.m"2)
100 120I I I • I1 I I 'l • ) I I • I
140
Figure 27: Crack growth rate versus stress intensity factor of F-82H modified on normalized (1075°C/30') state.
(From Ref. 16).
10
co
I 6~c01ocoO 4coD£
o
o
As Q.
0
o Ä
©
Maximum Hyc
Specimens
1 i
O
0
0
O
rogen
1
0
0°
Level
0
0
lor V
I
fi
1
¡
rgin
X
0
©A
i
403
F6NM
630
1500 600 700 800
Tempering Temperature (*C)
Figure 28: Hydrogen concentration measured in specimens after CBB test in 288DC water for 500 hours.
(From Ref. 29).
51
1BD0
SPUTTER TIME (sec)
3600 5400 7200
30 60 SO 120
ESTIMATED SPUTTER DEPTH (nm)
Figure 29: Concentration versus depth from the surface fracture macro-facets on the specimen irradiated at
410°C. The concentrations are the averaae values from seven areas on three facets. (From Ref. 39).
52
15
* I O
<er
o
JFMSÍB.7
i.OMcV, e" 5OO*C 3.ldpo
-100 0 KDO 200 3O0DISTANCE FROM GRAIN BOUNDARY (nm)
Figure 30: The solute concentration at and near a grain boundary in irradiated Japanese Ferritic /Martensitic
steels. The data points are the average of five to eight measurements at the same distance from the grain
boundary. The error bars connect the minimum and maximum values for the five to eight cases. (From Ref. 40).
1.0
C.5
— i
i °rl0
MATRIX
I ICo Ci
T« 7»' I ' I l.2MeV,e-
50C-C 2003C
eoj-601- 9* ¿*
• »CST-RRA0
ça o»-DO - 5 0 0 50
DISTANCE FROT/- ?=T-V¿THS< N
I n m )
00
Figure 31: The compositional profile near a MÔC type precipitate-matrix interface in Japanese
Ferritic/Martensitic steels before and after irradiation at-500°C-. (From Ref. 40).
53
LO
2LUuoo
LU0.5J-
¿7 MATRX S r m FROM! 1 PR£- RRAD
EZZ3 POST-RíAC L25**aV, s" 5OO-C 2Oopo
M2î
O1-
1
Me C
1
I.K
N¡ Si CrUOJ)
Ni Si Cr(JOJ)
Ni S¡ CrIIO.I)
Figure 32: The change in solute concentration at matrix in the vicinity of the precipitates in Japanese
Ferritic/Martensitic Steels by irradiation at 500°C. (From Ref. 40).
1.0
ÜJ
°0.5
AT PRECIPÍTATE 50nm FROM INTERFACEI ! PRE-lftRAO
EZZZ2 POST-RRAD l.25MeV , t* 5OCTC 2Ooro
MZ3 C6
t
1.55 Ms C M C
Ni Si Cr Ni Si CrUOOI) 1x05! (xO.li
Ni Si CrIiO.I)
Figure 33: Same as Figure 32 except at the precipitates in the vicinity of precipitate-matrix interface.
(From Ref. 40).
54
-tOO - « 0 -400 -200 0 200 -¿00 600 800from Groin Boynûory (rw>)
Figure 34: Concentration profile of Cr in Fe at near a grain boundary as a function of distance from the grain
boundary. (From Ref. 40).
55
10 20 30 40
)frad>ation dose, dpa
Figura 35: Dependences of grain boundary concentration of Si and Mn on the irradiation dose in (a) 9'7rCr-
2%Mn-l%W and (b) 12%Cr-6%Mn-l%W alloys. (From. Ref.43)
56