Corrosion of High Chromium Ferritic/Martensitic Steels in High Temperature Water

Informes Técnicos Ciemat 925 abril, 2000

Corrosion of High Chromium Fenitic1 Martensitic Steels in High Temperahire Water, A Literature Review

P. Fernández J. Lapeña F. Blazquez

Departamento de Fisión Nuclear

Toda correspondent en relación con este trabajo debe dirigirse al Servicio de Información y Documentación, Centro de Investigaciones Energéticas, Medioambientales y Tecnológicas, Ciudad Universitaria, 28040-MADRID, ESPAÑA.

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 -14226-1995 ISSN: 1135-9420 ÑIPO: 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 Steels in 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 temperature water 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 addition some investigations about radiation induced segregation (RIS) are shown in order to know the compositional changes 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 to the behaviour of these alloys to stress corrosion cracking seem to conclude that in these materials is necessary to optimize the temper temperature and to carry out the post-weld heat treatments properly in order to avoid stress corrosion cracking.

Corrosion en Agua a Alta Tempertaura de los Aceros Ferríticos/Martensíticos de Alto 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 la corrosión en agua a alta temperatura de los aceros ferríticos/martensíticos de alto contenido en cromo. El estudio se ha basado principalmente en valorar la respuesta de este tipo de materiales en corrosión generalizada, efectos de la irradiación en corrosión, corrosión-bajo tensión y corrosión bajo tensión asistida por 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 de estas 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 comportamiento 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 los tramientos 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 C Cr Mo Nb Si Ni Mn S Cu

2,25Cr-lMo 0,10 2.27 0,95 0.90 0,35 0,64 0,50 0,005 0,02

9Cr-lMo 0,10 9,14 1,05 - 0,76 0,23 0.51 0,002 0,01

9Cr-lMo 0,08 8,79 0,95 - 0,61 0,27 0,52 0,006 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 Heat C Mn P S Si Cr Al Ni Ti

405 2804-1 0,064 0,51 0,015 0,025 0,66 13,40 0,100 0,12 -

405 1781-1 0,054 0,53 0,014 0,004 0,56 13,55 0,130 0,41 -

405 1782-2A 0,056 0,53 0,017 0,006 0,55 13,40 0,110 0,36 -

410 2926-1 0,070 0.55 0,015 0,008 0,44 12,30 - 0,22 -

409 2725-1B 0,039 0,41 0,023 0,004 0,60 10,83 - 0,32 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 C Si Mn Cr Ni Mo Nb Ta V W N

LA7Ta 0,15 0,07 0,76 11,4 0,02 <0,01 <0,01 0,12 0,25 2,9 0,07

LA12Ta 0,16 0,03 0,80 9,8 0,02 <0,01 <0,01 0,10 0,27 0,85 0,04

LA12TaLC 0,09 0,03 1.01 8,9 0,02 <0,01 <0,01 0,09 0,39 0,76 0,02

FV448 0,10 0,46 0,86 10,7 0,65 0,60 0,26 0,14

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 Oxide Thickness by Optical Oxide Thickness by Weight Crystal Microscopy (fini) change Descaling (ujni) Depth of size and during the Total Inner Outer Total Inner Outer Corrosion Steel range (ujn) test (nig) oxide layer layer oxide layer layer (Uni)*

0.32 1.3 0.98 0.64 0.34 0.85 0.67 0.18 0.52 FV448 (0.12) (0.7) (0.18) (0.10) (0.11) (0,04) (0.03) (0,04) 80.05) n=8 0.3-4.0 n=17 n=17 n=!7 n=3 n=3 n=3 n=4

0.11 1.5 0.82 0.59 0.23 0.50 0.39 0.11 0.31 LA7Ta (0.09) (1.1) (0.15) (0.06) (0.11) (0.02) (0.04) (0.02) (0.02) n=S 0.3-7,8 n=20 n=20 n=20 n=3 n=3 n=3 n=3

0,32 1.2 1.04 0.73 0.31 0.95 0.81 0,14 0.57 LA12Ta (0.12) (0.6) (0.22) (0.10) (0.17) (0,06) (0.07) (0.02) (0.04) n=S 0.3-3.4 n=22 n=22 n=22 n=3 n=3 n=3 n=3

0.32 1,3 1.30 0,84 0.46 0,90 0.74 0.16 0.54 LA12TaLC (0,11) (0.7) (0.24) (0.14) (0,20) (0.06) (0,05) (0,02) (0,04) n=8 0.5-6.3 n=29 n=29 n=29 n=3 n=3 n=3 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. Oxide Thickness by Optical Oxide Tliickness by Weight Microscopy (pjm.) change Region of Descaling (uni)

during the Specimen Total Inner Outer Total Inner Outer Steel test (mg) Measured oxide layer layer oxide layer layer

LA12TaLC 1 1.34 1.17 0.77 0.40 1,45 0.96 0.49 (Unwelded) (0.02) Edge (0.12) (0.09) (0.10) (0.39) (0.21) (0.18) n=3 n=6 n=6 n=6 n=2 n=2 n=2

1.57 0.54 1,03 1.70 0.91 0.80 Edge (0.19) (0.06) (0.16) (0.14) (0.05) (0,10) LA12TaLC 2 n=S n=8 n=S n=2 n=2 n=2 -1.61 (Welded) (0,97) 1.43 0.45 0.97 1.16 0,76 0.40 n=3 Weld (0.69) (0.08) (0,63) (0,12) (0,07) (0,05) n=16 n=16 n=16 n=2 n=2 n=2

1.72 0,67 1.05 2,63 0.84 1.81 Edge (0.24) (0.08) (0,20) (0,33) (0.13) (0,22) n=9 n=9 n=9 n=2 n=2 n=2 -0.58 LA12TaLC 3 (0.25) (Welded) 2,20 1.62 0.58 1.43 0.89 0.54 n=4 Weld (1.37) (1,65) (0,54) (0,09) (0.11) (0.01) n=22 n=22 n=22 n=2 n=2 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. Oxide Thickness by Optical Oxide Thickness by Weight Crystal Microscopy (urn) Depth of change Descaling (urn) size and Corrosion Steel during the Total Inner Outer Total Inner Outer range (jim) (fini)* test (nig) oxide layer layer oxide layer layer

0.62 2.1 1.38 0.70 0.68 1.28 1.12 0,16 0.75 FV448 (0.19) (0.7) (0.16) (0.07) (0.13) (0.09) (0.09) (0.04) (0,09) n=8 1,1-4.8 n=20 n=20 n=20 n=3 n=3 n=3 n=3

0.47 2.0 1,36 0.58 0.78 0.99 0.82 0.17 0.57 LA7Ta (0.14) (0.5) (0.25) (0.11) (0.20) (0.009 (0.05) (0.05) (0.01) n=8 1.0-3.8 n=15 n=15 n=15 n=3 n=3 n=3 n=3

0.62 2.1 1.54 0.63 0.91 1.30 1.11 0.19 0.78 LA12Ta (0.18) (0.6) (0,18) (0,10) (0.24) (0.07) (0.03) (0,05) (0.02) n=8 1.1-4.0 n=15 n=15 n=15 n=3 n=3 n=3 n=3

0.62 0.1 1.52 0.62 0.89 1.37 1.20 0,16 0,80 LA12TaLC (0,18) (0.7) (0.18) (0,07) (0.15) (0.02) (0.04) (0.01) (0,05) n=8 0.6-5.2 n=16 n=16 n=16 n=3 n=3 n=3 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 Oxide Thickness by Optical Oxide Thickness by Weight Microscopy (um) change Region of Descaling (uni)

Steel during the Specimen Total Inner Outer Total Inner Outer test (mg) Measured oxide layer layer oxide layer layer

LA12TaLC 1 0,32 1.66 0.66 1,00 1.45 0.92 0.53 (Unwelded) (0.08) Edge (0.15) (0.08) (0.16) (0.28) (0.01) (0.29) n=2 n=10 n=l0 n=10 n=2 n=2 n=2

1.43 0.60 0.83 2.15 1.34 0.82 Edge (0.25) (0.06) (0.26) (0.13) (0.12) (0.01) -1.61 LA12TaLC 2 n=9 n=9 n=9 n=2 n=2 n=2 (0.97) (Welded)

n=3 1.77 0.65 1.12 1.57 0.94 0.56 Weld (0.76) (0.10) (0,68) (0,45) (0.27) (0,08) n=6 n=6 n=6 n=2 n=2 n=2

1.48 0.59 0.88 1.79 1.28 0,51 Edge (0.15) (0,08) (0,17) (0.04) (0.18) (0.23) LAllTaLC 3 -0,58 n=8 n=8 n=8 n=2 n=2 n=2 (Welded) (0,25)

1.99 1.16 0.83 1,27 0.91 0,36 n=4 Weld (0.93) (1.25) (0.55) (0.05) (0.18) (0.13) n=17 n=17 n=17 n=2 n=2 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 EA 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 Time to rupture (h) Time to rupture (h) Environment at room Notched* Tempered 250°C/2h Tempered 450°C/2h Temperature (HV 517) (HV 531)

Yes 186 210 Atmosphere Yes 310 7,2 No 2580 144

Yes 2,3 3% NaCI Yes 70 No 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.

Ref. Boiler Water Ref. Boiler Water 35000 h, 35000 h, 10% NaOH Cl" (100 ppm) Sample Condition 240 Mpa. Plast.def. 0,15a.v 4800 h, 0,9CTV 10600h, 0,9ay

As received NCll) NC NC NC

Tempered 650°C NC NC NC NC

Tempered 565°C NC SCC NC NC

Tempered 482°C NC SCC SCC 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 alu/. 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.

Aver, Current UTS CTy Surface T° (°C) tr(h) Elongation Type of Fracture Density, (ksi) (ksi) Appearence mA/cnr

25 57 105 68 20,52 Ductile cup & cone Black- + 0.008 50 56 107 71 20.16 Ductile General pH=2 cup & cone Corrosion +0,12

Eapp = 640 75 26 101 81 9.36 Brittle

mVH IGSCC&TGSCC Pitting +0,89 100 17 89,4 81,3 6.12 Brittle IGSCC&TGSCC Pitting +2,47

25 59 109.5 59.7 21,24 Ductile cup & cone Clean 0.0045 50 56 104 67 20.16 Ductile pH = 7 cup & cone Clean +0.009

Eapp = 540 75 24 97 83 8.64 Brittle Covered with m\'„ IGSCC&TGSCC Film +0,114 100 20 103 96.5 7.20 Brittle IGSCC&TGSCC Pitting +1.86

25 60 110 65 21.60 Ductile cup & cone Clean +0,004 50 53 108 76 19.08 Ductile Covered with pH =10 cup & cone Film +0.013

Eüpp = 540 75 34 95.6 74.4 12.24 Brittle mV,, IGSCC&TGSCC Pitting +0.86 100 22 103 66 7.92 Brittle IGSCC&TGSCC Pitting +0.543

Table 10: Stress Corrosion Test Result for type 403 Stainless Steel in 0,0JM Na2SO4, pH 2.7 and 10 at Various Temperatures. Starin rate = 10"-6 s"-1 .

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.

An effect of heat treatments on samples of martensitic steels has been reported by Tsubota et al(29) 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 C Si Mn P S Ni Cr Others

403 0,14 0,44 0,74 0,028 0,007 0,25 11,70 -

420J1 0,16 0,29 0,39 0,031 0,013 0,30 12,12 -

CA40 0,26 0,49 0,77 0,02 0,01 0,10 11,9 -

F6NM 0,033 0,36 0,65 0,02 0,013 3,87 12,74 Mo 0,50

431 0,18 0,82 0,80 0,036 0,011 2,42 16,98 -

630 0,04 0,29 0,74 0,027 0,003 4,06 15,71 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 HV30 Normalized at 1075°C/30' 405 Normalized at 1040°C/30' + Tempered at 750°C /I h air cooled. 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 showed

Q 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 REFERENCES

1. Harries, D.R., Dupouy, J.M. and Wu, C.H. Journal of Nuclear Materials. 133-134(1985)25.

2. Harries, D.R. Radiatt. Eff.101 (1986) 3.

3. Harries, D.R. , Butterworth, G.J. Hishinuma, A, and Wiffen, F.W. Journal of Nuclear Materials. 192-194(1992)92.

4. Proc. EEA Worshop (Working-Group -Annex H) on Ferritic /Martensitic Steels, Tokyo, October 1992.

5. Proc. IEA Workshop (Working Group- Annex II) on Low Activation Materials, Culham , April 1991, EUR Fu BRU/Xn-6/91-MATAI 16.

6. Mann, G.M.W. "The oxidation of iron base alloys containing less than 12%Cr in high temperature aqueous solutions". High Temperature High Pressure Electrochemistry in Aqueous Solutions, NACE, Houston, 1976, p 34-47.

7. Tomlinson L, et al. "Sodium heated steam generator tubes: Effect of heat flux on the deposition of magnetite from solution and corrosion of the underlying steel". Corrosion 41 (1985) p 257-264.

8. Tomlinson L, et al. "Deposition of magnetite from solution onto steam generator tubes surfaces and corrosion of the underlying steel". Proc. Conf. Water Chemistry of Nuclear Reactor Systems, Bournemouth, British Nuclear Energy Society, p.37, 1980.

9. "Ferritic Steels for Fast Reactors Steam Generators", Proceedings of the International Conference, Vols.l and 2 (S.F. Pugh and E.A. Little, eds., British Nuclear Energy Society, London (1978).

10. DeVan, J.H. and Gries, J.C. "Clinch River Breeder Reactor Environmental Effects-General Waterside Corrosion", Nuclear Technology. 28. 398-405 (1976).

11. Armijo. J.S., Krankota, J.L.. Spalaris. C.N.. Horst. K. M., and Tippits. F.E. "Materials Selection and Expected Performance in Near-Term LMFBR Steam Generators" . Proceedings of the International Conference on Fast Reactor Power Stations. pI89-203, Nuclear Energy Society, London (1974).

32 12. Broomfield J.P, et al. "Mechanism of corrosion and denting of ferritic steels in aqueous chloride at 300°C". Proc. Topical Conf on Ferritic Alloys for Use in Nuciera Energy Tech, Met Soc AME, Warrendale, PA, 1984, p 77-82.

13. Vaia A.R, et al. "Denting of steam generator Tubes in PWR plants". Materials Performance 19 (1980) p 9-24.

14. Petter E.C. , Mann G.M.W. "Oxidation of Mild Steel in High Temperature Aqueous Solutions". li! International Congress on Metallic Corrosion, Butterworths, London, p 416 (1961).

15. Ashmore C.B. and Large N.R. "Corrosion of Low Activation Martensitic Stainless Steels developed for Reactor Applications". AEA Fusion, Culham Laboratory, Abingdon, UK (AEA FUS 102).

16. Lapeña J. and Blázquez F. "Water corrosion of F-82H modified in simulated irradiation conditions by heat treatment". Proc. Conf. "9 International Conference on Fusion Reactor Materials" (ICFRM-9). October 10-15, Colorado Springs, EEUU. 1999. To be publish in Journal of Nuclear Materials.

17. Lapeña J. and Blázquez F. Ciemat Repon. To be published

18. Yamanouchi N. et al. "Accumaulation of engineering data for practical use of reduced activation ferritic steel: 8%Cr-2%W-0,2%V-0,04%Ta-Fe." Journal of Nuclear Materials. 191-194, p.822- 826(1992).

19. Tomlison L., Hurdus M.H., Ashmore C.B. and Silver Pi. B. Corrosion, Vol 41 n° 5, p 257-264 (1985).

20. Marsh G.P. et al. "The influence of radiation on the corrosion of stainless steel". Corr Sei 26 (1986) p 971-982.

21. Källström R. "The effects of irradiation and radiolysis on the aqueous corrosion of high chromium martensitic steels". Studvik Material AB. 1993 (STUDSVIK/M-93/106).

33 22. Gott K and Lind A. "Corrosion during the irradiation of Manet and Low Activation Alloys". 17th Symp on Fusion Techn, Rome, September 14-18. 1992, p30.

23. Truman J E. "Stress-Corrosion cracking of martensitic and ferritic Stainless Steels". Int Metals Rev (1981): 6 p 301-349.

24. Ozaki T and Ishikawa Y. "Stress Corrosion Cracking in nickel-bearing 13Cr Martensitic Stainless Steels in hot pure water". Corr Eng 38 (1989) p 165-174.

25. Ozaki T and Ishikawa Y. "Intergranular stress corrosion cracking and hydrogen embrittlement of martensitic stainless steels in high temperature, high purity water". Proc Int Conf on Stainless Steels, The Iron and Steel Institute of Japan, 1991, p 176-180.

26. Wilson I.L.W, Pement F.W and Aspden R.G . "Stress Corrosion studies on some stainless steels in elevated temperature aqueous environments". Corrosion 34 (1978) p 311-320.

27. Doig P., Chastell DJ., and Flewitt P.E. "The stress corrosion susceptibility of a quenched and tempered 12pct CrMoV martensitic stainless steel". Met Trans 13A (19829 p 913-921.

28. Bavarian B, Szklarska-Smialowska Z., and Macdonald DD. "Effect of temperature on the stress corrosion cracking of tempered type 403 martensitic stainless steel in sodium sulfate solution". Corrosion 38 (1982) p 604-608.

29. Tsubota M, Hattori K, and Okada T. "Characterization on long term aged martensitic steels". Proc Fifth Int Symp on Environmental Degradation of Materials in Nuclear Power Systems-Water Reactors, ANS, La Grange Park, EL, 1992, p 305-310.

30. Boler R et al. AEA Technology Report, AEA FUS 164, 1992.

31. Jacobs A.J. and Wazadlo G.P. "Irradiation-Assisted Stress Corrosion cracking as a factor in Nuclear Power Plant Aging". Proc. Int. Conf. of Nuclear Power Plant Aging, Availability Factor and Realiability Analysis, San Diego, CA, ASM, p 173.1985.

32. "Effects of Radiation on Materials". 14th Int. Symp, Andover, MA, June 1988, ASTM, Philadelphia.

33. "Radiation-Induced Changes in Microstructure, 13th Int. Symp. Parts I and II, Eds. F.A. Garner, 34 N.H. Packan, and A.S. Kumar,-STP 955 and 956, ASTM, Philadelphia, 1987.

34. C.C. Lin. "An Overwiev of radiation Chemistry in Reactor Coolants". Proc. Environmental Degradation of Materials in Nuclear Power Systems-Water Reactors, Monterey, Sept. 1985, ANS, 1986, p. 160-172.

35. Proc. of Water Chemistry of Nuclear Reactor Systems, 5, Bournemouth, October 23-27 1989, British Nuclear Energy Society, London, 1989.

36. Nystrand A-C. "The Effects of Irradiation on Stress Corrosion Cracking Sensitivity of a Martensitic Steel type 1.4914". Studsvik Material AB, 1994 (STUDS VK/M-94/117).

37. Nystrand A-C."The effect of Irradiation on Tensile properties and stress corrosion cracking sensitivity of martensitic steels". Studsvik Material AB, 1994 (STUDSVIK/M-94/18).

38. Morgan T.S., et al. "Interfacial segregation in fast reactor irradiated 12% chromium martensitic steel". Effects of Radiation on Materials: 15th Int Symp, ASTM STP 1125, ASTM, Philadelphia, 1992, p. 633-644.

39. Clausing R.E. et al. "Radiation-Induced segregation in HT-9 martensitic steel". Journal Nuclear Materials 141-143 (1986) p. 978-981.

40. Muroga T., Yamaguchi A., and Yoshida N. "Characteristics of radiation-induced solute segregation in candidate and model ferritic alloys". Effects of Radiation on Materials: 14th Int Symp, Vol 1, ASTM STP 1046, ASTM. Philadelphia, 1989, p 396-410.

41. Takahashi H., Ohnuki S., and Takeyama T. "Radiation-induced segregation at internal sinks in electron irradiated binary alloys". Journal of Nuclear Materials 103-104 (1981) p 1415-1419.

42. Ohnuki. S, Takahashi H, and Takeyama T. "Void swelling and segregation of solute in ion-irradiated ferritic steels". Journal of Nuclear Materials 103-104 (1981) p 1121-1125.

43. Kimura.A, Charlot.L.A. Gelles.D.S, Jones.R.H. "Dependence of grain boundary chemistry on the irradiation dose in low activation ferritics". Journal of Nuclear Materials 212-215 (1994) p.725-729.

35 44. Kimura.A, Charlot.L.A, Gelles.D.S, Baer. D.R, Jones.R.H. "Irradiation induced changes in the grain boundary chemistry of high-manganese low cativation martensitic steels". Journal of Nuclear 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 90

70 g E 60 SSO I ]

c O) 40 a Base metal (TIG) o Base metal (EB) 30 c Weldment metal (TIG) O Weldment metal (EB)

1000 2000 3000 4000 5000 leu {»ni Time (hours)

Figure 2: Metal loss versus time for 21/4 Cr and 9% Figura 4: Weight loss versus time for base and weld Cr ferritic steels. Average line is least squares fit to material of F-82H modified. (From Ref. 16 and 17) data. (From Ref.7).

10a

10'

«niai ai4X)SSj3 NOM «Jira uf í«l !¿Í1O'

rto*3a *o s

1O- - 1 i ? •

i Jim»- • un* ivi*t~

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- Figure 5: Time to rupture versus tempering- Fe;O3 on metal loss. (From Ref.7). temperature plots for three 13% Cr steels; smooth (a) AVT water chemistry; low dissolved O2 (<10 specimens tested in 37c NaCl solution at ambient ppb);4100h. temperature; initial stresses 695 or 1004 MNm"2. (b) AVT water chemistry; high dissolved O2 (56 (From Ref.23). ppb); 3500 h.

41 x O O7-/.C o O 26-/.C D O 32-/.C

O ,10'

2OO 3OO "JOO 5OO 6OO 7OO 2OC 3OO -aOO TEMPERING TEMPERATURE, *C TEMPERING

Figure 6: Time to rupture versus tempering Figure 8: Effect of tempering temperature on time to temperature plots for smooths specimens of 13% Cr rupture of 13% Cr steels stressed in a sulphide steels tested in boiling 3% NaCl solution; initial environment. Smooth samples stressed initially at 464 stress 464 MNni2. (From Ref. 23). Mnm":. (From Ref. 23).

O19--C I AC 1O5O-C OZB-Í.CI

0 naroness V Iiod ù corrosion rate —1 600< 0 n TH , =500 TO- L1»U C -3-Ou,- y- ci *J4 2^00 h— U Û < tr c. $ 27 ""300 û 0 10"' 2OC 3OQ &DC 7OC 200 10O¿ MNrrr' 695 MNn-;* 0 100 200 300 400 500 600700 TEMPERING TEMPERATURE (In.1, "C TEMPERATURE, *C

Figure 7: Effect of tempering temperature on time to Figure 9: Effects of tempering on hardness, impact rupture of notched samples of three 13% Cr steels strength, and corrosion rate in 10% HNO3 at 20=C of stressed in 1% HC1 at ambient temperature. (From hardened 0.28-13Cr steel. (From Ref. 23). Ref. 23).

42 I T 1 105O"/250-C condîfon tempérée, stress, MNrrf 3*/. Na Cl x 250 TOO! 10' o -áSO 10O4 1700- ù 550 695 cast O plain tensile speamens at 1OOOMN m'2 no

.adCnarpy nens at 22MNm 2

10 _L _ I.I -16 -12-6-4 0 4 8 12 TIME TO RUPTURE, h 4 vo -ve e- APPLIED CURRENT DENSITY, mA cm"2

Figure 10: Effect of strenght (UTS) on time to Figure 12: Effects of galvanostatic polarization on rupture of 13% Cr steels of variying carbon content time to rupture of a 13% Cr steel tempered at 250°C tempered at 250°C, 450 °C and 55OQC; smooth stressed in 3% NaCl solution at ambient temperature. specimens were tested in 3% NaCl solution at (From Ref. 23). ambient temperature. (From Ref. 23).

o \ tempered,°C stress,MNrr 1400 , X*X . OCA *r\r\j

ß:

x 0-ITÍ.C o Q22*IJZAC, 1050"C-2nst250'C D

- 200h

2 3 4 5 ¿ 3 4 5 10 10 10 10 10 )0 10 10 10 TIME TO RUPTURE , h TIME TO RUPTURE, n

Figure 11: Effect of strength (UTS) on time to Figure 13: Effect of initial aplied stress on time to rupture of 13% Cr steels of varying carbon content rupture of 13% Cr steels hardened and tempered at tempered at 250°, 450°. or 550°C; smooth specimens 250°C; corrodent used was 3% NaCl solution at were tested exposed to an industrial atmosphere. ambient temperature. (From Reñ 23) (From Ref. 23).

43 o x 01V/.C ] 1050*/250 condition o 0 lev.C I AC 1050'C • 2h at 250-C applied 1400 - D 028-/.C 'E x 2 D V o

UJ UJ

11O"J < 600- h- 2 O er O 200- V 10" <

3 u 10 10 10' 10= 2000 4 000 6000 800010000 TIME TO RUPTURE INITIAL STRESS INTENSITY, kg cm"3'J

Figure 14: Effect of initial applied stress (Smooth Figure 16: Effect of polarization on crack growth specimens) on time to rupture of 13%Cr steels rate of 0,21C-13%Cr steel hardened and tempered at hardened and tempered at 250°C, corrodent used 250°C; precracked specimens were tested in 3% was an industrial atmosphere. (From Ref. 23). NaCl solution at ambient temperature unpolarized, polarized anodically, or polarized cathodically. (From Ref. 23).

I 1 I 1050*/250* condition oocn symbols 'valid' tests E (with respect to specimen size) u dosed symbols 'invalid' tests 5000 01 (with respect to specimen size) o 0T/.C D 0-27.C V 0-3°/.C

1050°/250° . condition I/) NaO,% pH ZZ o 03 7 à 3-0 7

J x 100 7 10" 1 10 Kr 1O TIME TO RUPTURE, h 1 10 10' 1OJ TIME TO RUPTURE, h

Figure 15: Initial stress intensity versus time to Figure 17: Initial stress intensity versus time to rupture plots for thrree 13<7rCr steels hardened and rupture plots for 0,21C-13Cr steel hardened and tempered at 250=C; precracked samples tested in tempered at 250°C; precracked specimens were 37rNaCl solution at ambient temperature. (From Ref. tested in NaCl solutions of various concentrations at 23). room temperature. (From Ref. 23).

44 1O5O*/25O'C condition 3*/. NaCl precracked Cnarpy specimens at 2239-6 kg cm"3# 102

\ 10 1 \ J \ 0 "*

1 /

m-1 -1200 -800 -400 0 10' 1 10 102 1C TO" APPLIED POTENTIAL,mV(sce) TIME TO RUPTURE, h

Figure 18: Influence of applied potential Figure 19: Initial stress intensity -versus time to (Potenciostatic control) on time to rupture of rupture curves for ~ 0,39c carbon steels of 0-12% Cr precracked specimens of 0,21C-13Cr steel, hardened content tested in 1050°C and 200°C-250°C tempered and tempered at 250°C, in 3%NaCl solution at condition; precracked samples were tested in 3% ambient temperature. (From Ref. 23). NaCl solution at ambient temperature. (From Ref. 23)

45 10- n i i r i i r l$c «..i« ¿ SO v m/56h x unpoiari2ed ù cathodicaiiy polarized].,. -2 -10' 0 anodicaliy polarized J1DmAcm

10" 400 iCC tOO M • 500 ¡M ?0C 50C 60S !0C 01 234567 89 10 11 12 Cr, •/. Tempering !emp./"C x 5h

Figure 20: Influence of polarization (galvanostatic) Figure 21: Sumamry of tensile strenght and IGC on times to rupture of ~0,37c carbon steels of 0- depth for the 13% Cr steels with typical Ni content as 12%Cr content tested at initial stress intensity of a function tempering temperature. (From Ref. 24). 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).

46 too 5O0 60C 7DC Temperinc temp./^C x 5h

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).

• j 2&8"C our« «« lOÖCf L -- j e :scc ¡ j D • No SCC 1 o

500 -

O o to O

o -^> HC # V-N No 'ole V O Ä ' 10 - 0 O O CO e o! -V" 200 25C 300

Haroness / M-

Figure 23: The critical SCC susceptibility-hardness-IGC depth diagram. (From Ref. 24). 47 M I 0.06C-5~6NÍ-13Cr I. \ 10* - X K B X c CJ

200 250 300 350 ¿00 ¿50 200 250 300 350 ¿0C Hardness / Hv

Figure 24: The critical SCC susceptibility-hardness-IGC depth diagram for the typical C and Ni contents. (From Ref. 24)

48 (288 °C)

SCC o 620*0 r.o 4 Temper o '8

2 No SCC 6 • SCC 650 O no SCC 660 Temper

2 r- o fc 0 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 2000 ¡ m0— X 403 42CJ) £ 1500 - A D CA40 X 0 • 431 "c. c A 63C O 1000 je o 500 SC C c o og n As 0 500 60C 700 800 Tempering Temperature (C)

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

® with hydrogen a without hydrogen

ë. s 13 1E-7 O en o 2 ü

1E-8- 1 1 i ' • i ' ' ' ' ' ' ' ' i ' ' ' ' ' ' ' I I I • I I I 'l • ) I I • I 20 40 60 80 100 120 140 K (MPa.m"2)

Figure 27: Crack growth rate versus stress intensity factor of F-82H modified on normalized (1075°C/30') state. (From Ref. 16).

10 0 f i X 403 10 F6NM o 0 © c o Ä A 6 o O 0 ¡ 630 O ~cI 0 01 0 o c 0 0° o O 4 c o © D£ o Maximum Hyc rogen Level lor V rgin Specimens

1 i 1 I i 1 As Q. 500 600 700 800 Tempering Temperature (*C)

Figure 28: Hydrogen concentration measured in specimens after CBB test in 288DC water for 500 hours. (From Ref. 29).

51 SPUTTER TIME (sec) 1BD0 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 ÍB.7 JFMS i.OMcV, e" 5OO*C 3.ldpo 15

*IO

< er

-100 0 KDO 200 3O0 DISTANCE FROM GRAIN BOUNDARY (nm)

MATRIX

1.0

C.5

I I Co Ci

— i

i °lr0

T« 7» 'I 'I l.2MeV,e- 50C-C 2003C

eoj- 601- 9* ¿* • »CST-RRA0

ça o» -DO -50 0 50 00 DISTANCE FROT/- ?=T-V¿THS< N I nm )

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 ¿7 MATRX Srm FROM ! 1 PR£- RRAD EZZ3 POST-RíAC L25**aV, s" 5OO-C 2Oopo I.K M2î Me C

LO 1

2 1 LU u o o 0.5J- LU

O1- N¡ Si Cr Ni Si Cr Ni S¡ Cr UOJ) (JOJ) IIO.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).

AT PRECIPÍTATE 50nm FROM INTERFACE I ! PRE-lftRAO EZZZ2 POST-RRAD l.25MeV , t* 5OCTC 2Ooro 1.55 MZ3 C6 Ms C M C

1.0

ÜJ

°0.5 t

Ni Si Cr Ni Si Cr Ni Si Cr UOOI) 1x05! (xO.li IiO.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 -¿00 600 800 from 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)