ML20084Q676

Intergranular Corrosion Cracking of Type 304 Stainless Steel in Water Cooled Reactors

ML20084Q676
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Issue date:	04/07/1970
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4 Attaciunent A INTERGRANULAR CORROSION CRACKING OF TYPE 304 STAINLESS STEEL IN WATER-COOLED REACTORS C. F. Cheng s 5 Materials Science Division Argonne National Laboratory

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Tabic of Contents

1. Introduction

2. Instances of Service Failures 2.1 Boiling Water Reactors (BWR) 2.1.1 Vallecitos Boiling Water Reactors 2.1.2 Dresden-I Boiling Water Reactors 2.2 Pressurized Water Reactors (PWR) 2.2.1 PWR Test Loops . .'

2.2.2 Savannah River Reactor

3. Corrosion Studies on Intergranular Cracking 3.1 Impuritics in Aqueous Solutions 3.1.1 Nitric-Dichromate Solution 3.1.2 Water + 100 ppm FeC1 3

3.1.3 Water + 100 ppm Oxygen 3.2 Irradiation Embrittlement 3.2.1 Ilydrogen Embrittlement ,

4. Mechanisms of Intergranular Corrosion Cracking ,

4.1 Solute-Segregation Theory ,

4.2 Hydrogen Induc'ed Transformation

5. Conclusions References ,.

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INTERGRANULAR CORROSION CRACKING OF TYPE 304 STAINLESS STEEL IN WATER-COOLED REACTORSW l

C. F. Cheng

. Materials Science Division

. Argonne National Laboratory Argonne, Illinois 60439 ABSTRACT  ;

Nuclear pressure vessels, primary piping, core components,and fuel-cle-ment cladding are subject to cyclic thermal and mechanical stresses due to the high pressure and temperature of the coolant or of the' fuel. When Type 304 stainless steci is used as material for these parts, potential

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intergranularcorrosionfailuresduetothe}1ombinedactionof_corros'ionand~~

low-cyclic stress or high static stress need critical evaluation.

This paper reviews Type 304 stainless steel with respect to (a) the instances of service failure in water-cooled reactors, - (b) the various corrosion studies on intergranular cracking, and (c) the proposed mechanisms

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of intergranular corrosion cracking. .

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1. Introduction .

Nucicar pressure vessels, primary piping, core components, and fuel-clement cladding are subject to cyclic thermal and mechanic stresses due to the high pressure and temperature of the coolant or of the fuel.

When Type 304 stainicas steci is used as material for these parts, potential intergranular corrosion failures due to the combined action of corrosion and low-cyclic stress or high static stress need critical evaluation. Numerous instances of such intergranular corrosion cracking of austenitic stainicas steci (Typen 304 and 304L) have been observed .

In U. S. nuclear reactors. In boiling water reactors', service failures have been reported in a variety of components, including fuel-element cladding,( ' ) preheater pipes,( and vessel liners.(4} In pressurized water reactors, failures of fuel-clement cladding (5) and' outlet nozzles (6) have been reported. These failures have occurred in both metallurgical conditions; sensitized and solution annealed. The mode of cracking is exclusively intergranular, llowever, nonc_of the cracking instances is

. associated with cither chloride _or_causticaolution. This paper discusses austenitic stainicss steel with respect to:

(a) the instances'of the above-mentioned failures, (b) the various corrosion studies on intergranular cracking, and (c) the proposed mechanisms of intergranular corrosion.

2. Instances of Service Failure a .

Instances of intergranular corrosion cracking of austenitic stainicss stec1'(Types 304 and 304L) in water-cooled reactor service are briefly discussed below:

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'2.1 Boiling Water Reactors (BWR)

In boiling water reactors, the primary coolant is high-purity water I

at a temperature of about 286*C. The water normally operates at neutral l pil (N 7) with a resistivity of approximately 1 MG-cm total solids from . :

1 to 10 ppm, chloride of about 0.02 ppm, and oxygen from radiolytic decomposition up to 0.3 ppm in the water and from 10 to 20 ppm in the steam.(5) In-service intergranular fractures have been reported in two ,, ,

BWR. These reactors are located at Vallecitos, California and Dresden, ,

. Illinois.

2.1.1 Vallecitos Boiling Water Reactors (VBWR)

Type 304 stainicas steel tubing with outside diameter and wall thick-ness ranging from 0.914 to 1.07 cm (0.360 to 0.420 in.) and from 0.020 to 0.051 cm (0.508 to 0.020 in.), respectively, were tested as fuel-element cladding in VBWR. ( ) About 40 out of the 900 rods exhibited intergranular failures in 500-1000 hr. However, no carbide precipitate could be seen at the grain boundaribs by means of either light or cicctron microscopy. The cracking pa'tterns varied with'the initial con- .

dition of the tubing. Cracks in the cold-worked tubings were mostly .

longitudinal, whereas tihose in the annealed condition were predominantly-circumferential. A statistical analysis of the failure-rate data in-dicated the following:

(a) There was no significant difference in failure rate between solution annealed and cold-worked cladding.

(b) 'The failure rate of " collapsed cladding" was significantly higher.than for " free-standing cladding" (indicating the influence of cyclic stress).

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(c) Fuel rods that operate at high surface heat fluxes (e.g.,

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> 95 W/cm or 3'x 10 BTU /hr/f t ) showed an increased failure rate over fuel rods that operate at lower heat fluxes. ,

2.1.2 Dresden-I Boiling Water Reactor (Dresden-I BWR)

Type 304 stainless steel fuel-clement cladding exhibited longitudinal -

intergranular cracking in Dresden-I BWR.( The tubing used was ,

commercially annealed and then cold rolled to 1.242 cm (0.489 in.) ,

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, outsido diameter. The tubing rangcd from 0.025.to.0.043 in, in wall ._

2 thicknese. The cladding was subject to an average exposure of 95 W/cm (3 x 10 BTU /hr/f t ) peak heat flux. Since the cladding was "nonfree" sta'inding and the failure occurred only at the highest heat flux ragion, g it was concluded that stress was a strong contributor to the cracking mechanism. Fluence was not a factor, since simulated'ex-reactor tests showed similar failure phenomena for other 18 wt% Cr-8 wt% Ni austenitics as well. Furthermore, studies condt$cted by Hazelton et al.( ) on UO 2

fuel rods clad with Type 304 stainless steel'showed irra'diation produced ~ .

only slight structural changes of the cladding at similar fluence, but -

no increase in susceptibility to intergranular corrosion (as indicated by post-test exposure to acidified copper ' sulfate colution). Irradiation

~ exposures ranged from 1.7 x 10 0 to 9.8 x 10 20 c,2 (,y 3,y),

Preheater pipes of small diameter (6 in. or less) also suffered weld.

' defects'and leaking cracks in Dresden-I BWR'.( } 'The coolant leaks' that ocedrred in Type 304 stainless steel tubing; (Schedule 80) were

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t circumferential1y oriented cracks. The typical crack lengths were from 2 to 4 in. long. These cracks initiated at the inner surface of the pipe in the counterhored fit-up region adjacent to the welds.

Meta 11ographic examination revealed the mode of cracking was exclusively intergranular. The microstructure surrounding the cracks showed grain-boundary carbide precipitation during welding. The cracking was restricted to the heat-affected zones of the pipe base metal. The crack propagation initiated from a single gamma (y ) austenite phase of base metal of Type 304 stain 1 css steel and arrested at the two-phase auntenite plus delta ferrite (y + 6) weldment of Typc '308 stain 1 css steel.

However, no indication of cracking was observed in any of the large diameter (greater than 6 in.) piping. The small-diameter pipe was welded with a procedure that uses a high heat input and low number of passes, and the large-diameter piping was welded by using a low heat-input and multipic passes. The latter procedure produced less carbide precipitation than the former.

2.2 Pressurized Water Reactors (PWR)

In pressurized water reactors, the. primary coolant is also'high-purity water. The pH of the coolant, in some applications, is maintained in the range of f rom 9 to 10.5 by the addition of lithium or ammonium hydroxide.

In addition, hydrogen is added to suppress the oxygen formed from radio-lytic dissociation of water. Intergranular cladding failure of austenitic stainless steci in both sensitized and' solution annealed conditions has been reported in pressurized water reactor service.

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PWR Test Loops ,

One experiment conducted in Engineering Test Reactor (ETR) loop was prematurely terminated as a result of release of excessive fission product.(10) It was determined that a pump failure and subsequent loss of coolant were responsibic for the severe thermal conditions which cracked a fuel-rod cladding to produce the " fission break". The fuel .

rods had operated for 76 days in 600*F lithiated water at pH 10 and at-peak surface heat flux of 93 W/cm (296,000 BTU /hr/ft ). ,

The UO 2 Pellets were clad in Type 304L stainless steel tubing that had a diameter of 0.820 cm (0.332 in.). The stretch-formed fuel rod with 0.038-cm (15 mil) wall tubing suffered intergranular cracking, while the free standing fuel rod with' O.053-cm (21 mil) wall was only dimensionally distorted.

The severe thermal condition also changed the cladding from a solution annealed to sensitized structure. The cladding failure of the stretch-formed fuel rod was attributed to stress rupture in conjunction with hydrogen under the severe thermal conditions. The hydrogen was generated by the reaction of the UO fuel and water.

2 Knolls Atomic Power Laboratory (KAPL) reported intergran'lar cracking

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u of vacuum-annealed Type 304L stainicss steci cladding on poison elements used in test pressurized water reactors.(5,ll) The failure occurred during irradiation in 600*F ammoniated water at pH 10 when the cladding was stressed in excess of the yield point. The hydrogen content of the irradiated cladding increased from 4-6 ppm to 8-14 ppm in 700 effective full-power hours. Stress rupture cracking aided by hydrogen was indicated.

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- n-9 2.2.2 Savannah River Reactor Intergranular stress corrosion cracking was detected in two Type 304 stainless steel nozzles in the high-purity heavy-water moderator of the Savannah River Reactor Plant. '

The cracks propagated from surface intergranular attack at sensitized grain boundaries. The attack was the

. crilt of acid (HF-11N0 ) iP ckling during initial fabrication of the 3

nozzles. Cracking was attributed to chlorides although the bulk water contained only 0.006 ppm chloride. Fluoride contamination is more likely.

Recent studies reported that fluoride ions with a concentration below 1 ppm in bulk water can result in intergranular cracking at the crevices of sensitized Type 304 stainless steel (see sections 3.1.3 and 3.1.4).

3. Corrosion Studies on Intergranular cracking The service failures noted in Section (2.) are associated with one of die following two conditions:

(a) impurities in aqueous solution, and (b) embrittlement and irradiation effects.

We shall review laboratory studies on intergranular corrosion cracking under these conditions.

3.1 Impurities in Aqueous Solutions 3.1.1 Nitric-Dichromate Solution Sensitization is not a prerequisite for intergranular corrosion of Type 304 stainless steel. Nonsensitized steel will also corrode inter-granularly when exposed to certain chemical environments. Accelerated a

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corrosion of nonsensitized stain 1 css steel during exposure to boiling nitric acid was reported. as early as 1952 by DeLong(10 and later by Shirley and Truman.(15) These workers attributed this accelerated corrosion to the presence of chromium ions that had been dissolved from the test specimens. Others I0' have shown that dissolved chromium in nitric acid can accelerate intergranular attack only when the chromium is at the highest oxidation state, Cr+6 .

It was not until 1959, however, that Streichcr(18) reported that the accelerated corrosion of nonsensitized stainicas steel in nitric-dichromate solutions resulted in preferential grain-boundary attack. In addition to Cr+6, the following ions have been found to activate interg'ranular corrosion of nonsensitized stainless steci in n'itric acid solutions:

Mn (Mn0 -1), Ce , and Fe+ ( ' ) . The kinetics of intergranular 4

corrosion of nonsensitized stainless steel in nitric-dichromate solutions have been studied with respect to several environmental variables.

Various workers have shown that the' rate of intergranular attack increases

[ with corresponding increases in HNO concentration, C b concentration, 3 ,

solution temperature or specimen stress 1cvel.

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f Metallurgical variables that have been studied to date _ include (a) grain size of austenitic stainless steel, (b) cold work, (c) surface

! treatment, (d) alloy composition, and (c) heat treatment in nitric-

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I dichromate solutions. Coriou, et al. (1 ) and Armijo( } have studied the f effects of grain size and cold work. They reported that the rate of 1

intergranular penetration increases as the grain size increases, and decreases with an increase of cold work. The~ effects of these variables l

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are small; for example, a tenfold increase in grain size will result in only a twofold increase'in intergranular corrosion rate. Similarly, a 50 percent reduction in thickness (by rolling).will produce only a twofold decrease in corrosion rate. Warzee and Bergc(21) have found that superficial cold work reduces the rate of intergranular corrosion.

This surface working is effective in reducing the corrosion of both -

sensitized and nonsensitized alloys. Armijo( ) found that high-purity

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- Fe-Cr-Ni alloys (14 wt% Cr-14 _wt% Ni-ba[ lance Fe) )are immunc to l'ntergranular corrosion. The addition of impuritics such as C, N, 0, Mn and S does not promote intergranular corrosion susceptibility of a nonsensitized high-purity Fe-Cr-Ni alloy. However, the addition of Si and P will promote severe susceptibility. Furthermore, the removal of Si and P from austenitic alloys of commercial purity decreases susceptibility to inter-granular corrosion by a factor of six to eight. The effect of heat treatment has been extensively evaluated by Aust( } to support the solute segregation theory (see Section 4).' Uc reported that Type'304 stainless steelspecimensquenchedfromhightemperatures(1300*Cfweremore .

susceptibic to intergranular corrosion than specimens quenched from .

temperatures between 1200 and 900*C, Also, specimens stabilized at 800-900*C were more resistant to attack than solution-annealed specimens quenched at 1050*C.

+

3.1.2 ~ Water + 100 ppm FcCl 3

Pickett, et al.(24) have reported several solutions that produce intergranular cracking of nonsensitized stainicas steel. These included water solutions of FeC1 , FcCl " 2, Cr0 3, and NaOH at a' temperature of 3 2

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I e o l (d C i 345'C. In these solutions cracking has been observed only on stressed specimens. It was found that the type of intergranular cracking produced in 100 ppm iron chloride water closely resembled the in-reactor cracking in appearance, and occurred over the same range of temperatures.

3.1.3 Fluoride Contaminated Water Recent laboratory testa ( have shown that Type 304 stainicss steel

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is susceptibic to intergranular fluoride icn (F ) attack in neutral or alkaline water solutions. Extensive metallographic study of the corroded specimens demonstrated that F~ attack was purely intergranular and was limited to specimens with heavily sensitized structure. Neither the solution-annealed material nor the heat effected zone adjacent to weld-ment produced any trace of F~ attack under the same testing conditions.

The attack was consistently less severe on sensitized material of finer grain size and/or lower carbon content. Specimens of other materials

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appeared to be completely resistant to F attack after a sensitizing heat treatment at 1150*F (621*C) for from 10 to 24 hr. These include Inconel 600, Incoloy 800, fine-grain Type 304 stainless steel, fine-grain Type 347 stainless steel, and austeno-ferritic wcld deposits or castings.

Attack occurred at room' temperature and was more rapid at 82*C (180*F).

Data are not yet sufficient to define the F~ concentration necessary for attack to occur, llowever, in a number of tests at 180*F, concentrations Icss than 1 ppm F~ in bulk solution caused slight attack in crevices.

A very high concentration (4 wt% NaF) was not corrosive. The rate

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and mode of F attack was not affected by the amount of Cl in the corrosive solution over a range from about 15 ppm to less than 30 ppb. With respect se

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i to possible mechanistic interpretation, the following patterns are significant:

1. Cracking occurs at room temperature.

2. Wide changes in the pil have no effect (pil range tested 4.70-10.45).

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3. Very high F concentrations stop cracking. .

4. Intergranular attack without stress occurs with air formed oxide; whereas, intergranular attack on oxide free surface occura only eith stress (e.acpt in crevices).

3.1.4 Water + 100 ppm Oxygen General Electric Co. at Vallecitos used a bellow-loading system to study the intergranular cracking on austenitic~ stainless steel in water containing 100 ppm oxygen at a temperature of 289'c (554*F),(25,26) The tensile specimen was stressed by adjusting the load to 0.2% offset yield strength (at 289'C). The materials used in this study include Type 304 stainless steci (0.07 C) Type 304L stainless steel (0.02 C), 304 + Si (0.07 C, 4.1 Si), nickel doped 304 + Si (0.06 C, 4.0 Si, 14.3 Ni), ,

U. S. Steel 18-18-2 (0.06 C, 1.92 Si), Uranus S (0.02 C,'4.2 Si 13.9 Ni, .

17.6 Cr) and Uranus-50.p0.05 C, 0.51 Si, 7.9 Ni, 20.3 Cr). Three heat treatments were employed:

i) Solution annealed at 1095'C (2000*F) and water quenched

11) Sensitized, 24 hr at 595'c (1090*F) in vacuo and furnace cooled iii) s welded.

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O E l No evidence of localized attack was found after 300 hr of exposure, when tested in the solution-a'nnealed condition. In the sensitized condition, only those austenitic alloys (e.g. , Type 304 stainless steel, nickel doped 304 + Si and U. S. Steel 18-18-2) with continuous shect-type .

carbides in the grain boundaries exhibited cracking in 17-23 hr. Other austenitic alloys (e.g., Type 304L stainicas steci and Uranus S) that were relatively free of continuous shect-type carbides in the grain boundaries or duplex structure alloys (austenite + ferrite - c.g. , Type 304 stainicss steci + Si and Uranus-50) with no carbide precipitation in the austenite grain boundaries showed no failure in the 300-hr test.

~ ._?- ) To explore further the intergranular stress-corrosion behavior of sensitized austenitic alloys of varying carbon content, a series of carbon-doped, high-purity alloy (18 wt% Cr-9 wt%.Ni),_with from,0.02 to 0.15 wt% carbon, was likewise tested in the sensitized condition. The time-to-failure decreased rapidly when the carbon content exceeded 0.03 wt%.

In addition, sensitization behavior of welded heavy plate section of Types 304 and 304L stainicss steel was evaluated. The welds were made .

on 0.5-in.-thick plates (6 by 6 in.) by using a shicided argon-are process employing 3 passes at high heat input with Type 308 stainicas steel filler material. Tensile specimens were then cut from the plate welds that have the wcld bead in the middle of the gage section.

Duplicate specimens of a Type 304 stainicss steel wcld failed in 40 and 45 hr, and corresponding specimens of a Type 304L stainicss steel wcld survived in similar stress-corrosion tests for the entire period of 300 hr. Meta 11ographic examination revealed the intergranular cracking

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occurred through a region of heavy carbides 0.25 in. from the Type 304 stainicas steel wcld bead. The Type 308 stainless steel weld displayed no cracking nor was carbido precipitation visibic, as in the case of the Type 304L stainicas steci weld head. ,i The following conclusion may be drawn from the above work:

1. Solution-anncaled material with no plastic deformation, other .

than that associated with stressing to. yield point, showed no evidence of intergranular stress corrosion after exposure to 289'C water + 100 ppm 02 '

2. Sensitized austenitic materials (C < 0.03 wt%) do not exhibit intergranular stress corrosion in the above environment.

3. Addition of Si to the austenitic alloy'is of no benefit, if the carbon content is sufficiently large to cause carbido precipita-tion during sensitization.

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4. Duplex alloys (ferrite + austenite) appear to be immune to intergranular stress corrosion even after a heat treatment that would normally result in sensitization in austenitic materials. .,

'i j This immunity is independent of the' nature of the ferrite i

k stabilizer employed.

Recently, Wilde( ) cvaluated kinetics of crack initiation of sensitized Type 304 stainicas steel. lie observed:

,. (a)- the induction time (T N) f r crac ng in reases markedly when 1

applied stress decreases from 30,000 to 20,000 lb/sq. in., and l (b) cracking appears to nucicate in the plastically deformed area i-of the surface.

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In addition, surface pickling with HNO 3-HF on Type 304 stainicss steel af ter sensitization (at 621*C (1150*F) for 12 hr) lowers the cracking period from 5 to 10 times compared with corresponding sensitized specimens without llNO3 -HF pickling. It should be pointed out that this ,

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surface treatment may be related to F at tack. (13) However, crevices produced at the grain boundary by pickling is not a likely factor in

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the absence of F . For exampic Armijo(25) pickled solution-annealed specimens of Type 304 stainicss steel in a boiling nitrate-chromate solution to produce deep (40.015 cm), uniform, intergranular corrosion at the specimen surface., The specimen was then washed in boiling water and no stress corrosion failure occurred in the 289'c water + 100 ppm oxygen test.

3.2 Irradiation Embrittlement Hydrogen and helium, which can embrittic austenitic stainless steel, are known to be present in the nuc1 car reactors. Hydrogen is generated from radiolytic dissociation of water in thermal flux reactors. Hydrogen is also produced by transmutations in (n,p) reactions in the matrix .

elements (Fe, Ni, and Cr) of austenitic stainicss stccl. Similarly, helium .

is formed by transmutations in residual boron and in the (n,a) reaction

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in the matrix clements. Helium is currently considered a major cause.of irradiation embrittlement in fast breeder reactors,_but_not_in_ water-cooled reactors. In the latter case, the operating temperature is too low for grain-boundary shcaring. For this reason only hydrogen embrittlement will be discussed.

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3.2.1 Ilydronen Embrittlement -

9 Whiteman and Trorairo( '

showed that thin sections of cathodically hydrogenated austenitic stainless steci suffered a significant loss of ductility in analogy with well-characterized hydrogen embrittlement of ferritic steel. A similar result was observed by Cheng( } and by Vennett and Ansell(30) on Type 304L stainless steel when tested in a high-pressure hydrogen atmosphere. .

4. Mechanisms of Intergranular Corrosion Cracking In virtually all known cases of intergranular brittleness. it can be shown that the brittleness is associated with the presence of some-impurity. Presumably, the impurity is present,in the narrow grain-boundary interface in much higher concentration than the mean composition of the metal would indicate.(31) There has been for many years qualitative evidence for r.egregation of solutes to grain boundaries. -This evidence has been reviewed by Inman and Tipler( 2) in 1963, and by Westtrook( }

in 1964.

.In the past year or so, two new quantitative -techniques.have emerged to study the structure and composition of the grain boundaries. Marcus and Palmberg( } used Auger Electron Spectrometry to study antimony

- segregation at the intergranular fracture surface of embrittled AISI 3340.

Brenner et al.(31,35) are using field 1 ionization atom probes to

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explore atomic structure at a grain boundary and to identify thefindividual-l atoms in that structure.

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4.1 Solute-Senrenation Theory In nitric-dichromate' solution, Aust( ) postulated that chemical con-centration differences exist between grains and grain boundarics, that is, impurity segregation at boundaries. These chemical differences provide the driving force for localized attack. It is this impurity segregation that can Icad to accelerated dissolution of grain boundaries when the alloy is exposed to a suitabic corrodant. Thus, a rather simpic model ,

- for intergranular corrosion is shown schematically in Fig. 1. It was proposed ( } that the intergranular corrosion of austenitic stainicss steci is associated with the presence of continuous grain-boundary paths of either soluto-segregated (e.g. , phosphorus) or second-phase (carbide) regions. For Type 304 stainless steci in the solution-treated condition, that is, quenched from 1050' to about 1300*C, the solute ' segregation at the grain boundaries provides a narrow continuous path for attack by the corrosion solution, as shown in Fig.1(a). If the austenitic stain-Icss steci is quenched from 600*C (or from the sensitization heat-treatment range of 500 to 700*C), continuous carbides or second ph'ase. regions are ,

formed at the grain boundaries, Fig.1(b), thereby producing a wide ,

continuous path for cla'ssical intergranular corrosion in the sensitized condition. In material heat treated and quenched from 800 to 900*C,

-isolated carbides may incorporate solute impurities drained from the grain-boundary region.- This latter heat treatment (i.e., quench from 800 to 900*C) should then reduce'the intergranular corrosion rate, since a discontinuous penetration path for the corrosion medium is now present at the grain boundaries. In addition the reduced corrosion rate'should 1

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remain unchanged even after subsequent 150 hr sensitization at 650'C.

The reason is that carbide precipitation at 850*C should reduce the carbon supersaturation sufficiently to prevent further nucleation of the carbide at-the lower temperature of 650*C, despite the decrease in the solubility ,

of the carbon at the lower temperature. Also on the basis of this model, a low corrosion rate would be expected for a high-purity material having relatively cican or segregation-free grain boundaries, as shown in ,_

Fig. 1(d) . The results of the Aust et al.( 0' ) investigation-in

- nitric-chromate solution c1carly supported this model. 5 According to the chromium-depletion theory, the beneficial effect of a stabilization heat treatment (800 to 900*C) on corrosion of 'austenitic stainicas stcc1 is due to the diffusion of chromium into the chromium-depleted areas adjacent to thc precipitated carbides. Ilowever, as was -

shown previously when Type 304 stainless steci is slowly cooled from 1300 to 900*C, held for 1 hr at 900*C and quenched, a high rate of intergranular attack is still observed. In other words,.at the same annealing temperature and time where chromium diffusion is thought to -

remove chromium-depicted regions, poor corrosion resistance is found, -

which is due to the observed continuous carbides present' at the grain boundaries.(39) This result is in agreement with the Aust model( )

but not with the chromium-depiction theory. _

One of the important problems remaining was to determine the identity and state of the solute segregation that was involved in the inter-granular corrosion of solut' ion-treated and quenched austenitic stainless steel. .The fact that some type of solute segregation is responsible is i

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indicated by the work of Chaudron(40) who found that high-purity stain- '

icas steel is immune to grain-boundary corrosion in boiling nitric-dichromate solutions. The recent results by llanneman and Aust( } have suggestod that the mechanism of nonequilibrium solute segregation induced by excess vacancies may also be responsible for intergranular corrosion of solution-treated and quenched austenitic stain 1 css steel. For i

exampic, both excess grain-boundary hardening (Fig. 2) and intergranular corrosion are observed in solution-treated Type 304 stainless steel.

llowever, when a high purity austenitic stainlcos steel containing no excess grain-boundary hardening (01} is tested, there is no evidence of intergranular corrosion.( In addition, it is found( } that a solute such as phosphorus (0.01-0.1 wt%) promotes susceptibility to

. intergranular corrosion of solution-treated, high-purity'austenitic stainicss steel in nitric-dichromate solutions, whereas carbon (up to 0.1 wt%) has litti,e effect (Fig. 3). The boundary hardening is found with carbon additions, except in the highest carbon alloy tested.

Electron microscopy examination revealed the presence of carbide at the grain boundary.( }

Finally, the solute'-segregation theory of intergranular corrosion of austenitic stainless steel is usef'ul in analyzing the service failures in water-cooled reactors (Section 2) and in aqueous solution tests (Section 3). The only differences among these corrodents are rate of intergranular attack. In the case of duplex alloys (ferrit'e + austenite)

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in water '+ 100_ ppm oxygen tests (Section 3.1.3), . it has been reported (42) that the carbon precipitates at the ferrite /austenite boundaries an'd

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austenite boundaries. Thus, the latter boundaries probably have

" stabilized" or " segregated" structures rather than " sensitized" structures.

4.2 Ilydronen; Induced Transformation The mechanism of hydrogen embrittlement in austenitic stainless steel due to a hydrogen-induced transformation was recently clucidated by Whiteman( ) and by llolzworth and Louthan. ) Cathodic hydrogenation of austenitic stainless steel led to partial transformation of the y (fcc) austenitic lattice to the martensitic phases, a' (bec) and c (hep) .

The latter phases were the same as those formed by cold working hydrogen f ree austenite at very low temperatures. Whitenun further showed that

, absorption of hydrogen led both to internal strain, induced by an expansion of the y (fcc) lattice, and to a significant lowering of the stacking-fault energy (related to the activation energy barrier inhibiting phase transformation). As a consequence, the martensitic transformation was induced even in normally stable Type 310 stainless steel (20% Ni, 25% Cr), which normally does not undergo a strain-induced martensitic transformation even at -4 K. Benson et al. (44) have further shown that brittle intergranular fracture occurs along martensitic platelets formed in the presence of hydrogen atmosphere. These results indicate that the presence of hydrogen causes embrittlement and initiates microcracks at the strain-induced martensitic platelets in 'the steel.

af

4

,0

,8.,

()

1

()

~ 20 -

Conclusions Nuclear pressure vessels, primary piping, core components, and fuel-clement cladding are subject to cyclic thermal and mechanical stresses due to the high pressure and temperature of the coolant or of the fuel.

When Type 304 stainless steel is used as material for these parts in water-cooled reactors, instances of intergranular corrosion failure due to the combined action of corrosion and low-cyclic stress or high static stress have been occasionally observed.

Sensitization is not a prerequisite for intergranular corrosion of Type 304 stainless stccl. Under certain conditions, solution-annealed Type 304 stain 1 css steel will behave likewise. This is evidenced by the service failures in water-cooled reactors and by test results in laboratory corrosion studies.

Based on the solute-segregation theory Aust( postulated that intergranular corrosion of austenitic stainless steel is associated with localized attack on the impurities segregated at the grain boundaries.

This is the reason that a sensitized condition, which has a wide continuous path of either solute-segregated (e.g. phosphorous)' or second phasc~

(carbide) regions, is more susceptible to intergranular corrosion cracking than a solution-annealed condition with a narrow contiduous path of segregates. Similarly, a stabilized condition with a discontinuous path of isolated carbide is expected to be even less susceptible to inter-granular corrosion cracting. A high-purity material that has relatively clean or segregation-free grain boundaries is immune to such cracking.

M y.-- A

' 0^

[')': /9 l )

%  %.J }

l l

l The presence of hydrogen can cause embrittlement and also initiates microcracks at the strain-induced martensitic platelets in Type 304 stainless steel.

The ir.tcrgranular corrosion failures of Type 304 stainless steel in water-cooled reactors may be related to (a) selective chemical attack of impuritics segregated at grain boundaries or (b) hydrogen embrittlement,.

or both.

k 4

+

k e

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4

• d se

s . A Q .

- P.cferences-(1) W.11. Arlt and S. R. Vanderberg, Gen. Elec. Co. Report GEAP-43,60 (1963).

(2) C. N. Spalaris, Nucleonics, 21, p. 41 (1963).

.(3) J . P. liiggins, Nuclear hews, Nov. , p. 37 (1968). -

(4) N. Inlai, C. R. Sut. ton, E. A. Winunc, and R. F. Jones, Argonne National 1.aboratory Report' ANL-7117 (1965) .

. (5) T. J..Pashos, TID-8540 (1964).

(6) S. R. Rideout, Savanah River Laboratory Report DPSPU 62-30-26 (1963).

(7) R. N. Duncan et al. , Nuclear Applications, Vol.1, No. 5, p. 413.(1965).

(8) S. Naymark and C. N. Sp' alaris, International Geneva Conference on

e. .

Peaccrul use of Atomic Energy, Vol.11, p. 432 (1964).

. (9) W. S. IIazelton, C. J. Kubit and D. R. McClintock, Trans. Am. Nucl.

Soc. Vol. 6, No. 2, p . 388 (Nov . 1963) .

(10) D. C. Bullington, USAEC Report YAEC-182,'(Jan. 1962). -

(11) C. F. Cheng, Corronion, p. 341 (Nov.1964). & unpublished data.

(12) W. E. Eerry, Reactor Materials, Vol. 7, No. 1, (1964).

(13) C. T. Ward, D. L. Mathis and R. W. Stachic, Corrosion, p. 394~ (1969) .-

(14) W.11. DeLong, ASTM Spec. Tech. Pub. No. 93, p. 211 (1949) . *

(15) 11. T. Shirley and J. E. Trumann, Jn. Iron and Stoc"1 Inst. , Vol.171,

p. 304 (1952).

.(16) .A. 11. McIntosh, Chemistry and Industry, p... 667 (1957).

(17) J . = E. Truman, Jn. Appl. Chem. , Vol. ' 4, p. 273 (1954)'. -. ...

(16) M. A. Stretcher, J. Electrochem. Soc. Vol.106, p.161- (1959) . - .- -.c-D P

r-v -( p

.- r

-. I l

l

[

l

[

Refere_nces

. (19) 11. Corf ou et al. , Colloque 'de Metallurgy IV, p. 75 (1960).

(20) ~J . S. Armijo, Corrosion, p. 235 (1965).

(21) M. Warzec and P. Berge, Comptes Hendons, Vol. 259, p.144 (1964).

(22) J. S. Armijo, Corrosion, p. 24 (1968).

(23) K. T. Aust, Trans. of Met. Soc. of AIME, Vol. 245, p. 2117 (Oct.1969).

(24) A. E. Pickett et al. , Nuclear Applications, Vol.1, No. 5, p. 453 (1965). .

(25) J. S. Armijo, Corrosion, p. 319 (1968).

(26) Jersey Central Power and Light Co., USAEC Doc,ket No. 50-219, Amendment No. 37, Reactor Vessel Repair Program - Oyster Crcok Nuc1 car Power Plant No. 1 (April 1968).

.(27) R. E. Wilde, Corrosion, p. 359 (1969).

(28)- M. D. Whiteman. PhD. Thesis. Case Inst. of Tech. (1965).

(29) M. B. Whiteman and A. R. Trorano, Corrosion, p. 53 (1965).

(30)

R. M. Vennett and G. S. Ansell, Trans ASM, Vol. 60, p. 242 (1967).

(31) J. R. Low, Trans. TMS-AIME, Vol. 245, p. 2481 (1969) .

(32)

M. C. Inman and 11. R. Tipler, Met. Review, Vol. 8, p.105 (1963).

(33) J.11. Westbrook, Met. Review, Vol. 9, p. 415 (1964) .

(34) 11. Marcus and P. W. Palmberg, Trans. TMS, Vol. 245, p.1664.

(35) S. S. Brenner and J. Y. McKinney, Appl. Phys. Letters, Vol. 13, _

• .j

p. 29 (1968).

(36) ' K. T. Aust, TMS-AIME, Vol. 245, p. 2117 (1969) .

. (37) K. T. Aust, J. S. Armijo, and J.11. Westbrook, ASTM Trans. Quart. ,

Vol. 59, p. 544 (1966) . .

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p. - - .-e t

~ ,

( .

I e t r"3i  !

n, .e e

U V (38) R. E. llanneman and K. T. Aust, Scripta Met. Vol. 2, p. 235 (1968).

(39) K. T. Aust, J. S. Armijo, E. F. Koch and J. II. Westbrook, ASM Trans. Vol. 60, p. 360 (1967).

(40) G. Chaudron, EURAEC-796, Quart. Report 6, Oct.-Dec. (1963).

(41) K. T. Aust, J. S. Armijo, E. F. Koch and J. II. Westbrook, ASM Trans.

Quart. Vol. 61, p. 270 (1968).

(42) Metals llandbook,1, Properties and Selection of Metals, 8th Edition,

p. 435 (1961).

(43) M. L. Ilozworth and M. R. Louthan Jr., Corrosion, p. 110 (1968).

(44) R. B. Benson Jr. , R. K. Dann and L. W. Roberts, Jr. , Trans. Met.

Soc. AIME 242, p. 2199 (1968).

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1 i

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1

. y,-

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f .

i (c) SEGREGATED (b) SENSITIZED i . .

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l (c) STA3!t! ZED (c) ctE Ax - ,

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, Fig. 1. Schematic illustration of t ic model proposed for 1

intergranular corrosion. 41

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. Fig. 2. liardness-distance profile near a grain

[ boundary in solution-treated Type 304

• austenitic stainless steel (1-g load, l

5-sec loading time). (45) t, l

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Fig. 3. Ef f cet of phosphorus or carbon additions on the .

corrosion resistance of solution-treated high-

'j purity austenitic stainless steci (14 wt% Cr,

, 14 wt% Ni, balance Fe) 2 n boiling nitric 2 dichromate solution.

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