STRESS CORROSION CRACKING OF STAINLESS STEELS TESTED …envdeg2015.org/final-proceedings/ENVDEG/papers/ENVDEG126.pdf · (1) AREVA NP Technical Center, 30 Boulevard de l'Industrie

17th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors August 9-13, 2015, Ottawa, Ontario, Canada

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STRESS CORROSION CRACKING OF STAINLESS STEELS TESTED BY DYNAMIC LOADING IN OXIDIZING AND REDUCING PWR PRIMARY ENVIRONMENT

Nicolas Huin(1), Olivier Calonne(1), Steffen Berger(2), Bastian Devrient(2), Renate Kilian(2), Lionel Fournier(3) and Antoine Marion(3)

(1) AREVA NP Technical Center, 30 Boulevard de l'Industrie - Espace Magenta - BP 181 71205 Le Creusot Cedex France (2) AREVA GmbH Technical Center, Paul-Gossen-Str. 100, 91052 Erlangen, Germany (3)AREVA NP Tour AREVA 92084 Paris La Défense France

ABSRACT

Austenitic stainless steels (SS) are widely used for LWR component manufacturing. The field experience of Type 304, 316, 316 Ti and 347 austenitic stainless steels and their L grade in PWR primary circuits has been generally outstanding. However, a recent review made by EPRI of all known SCC of SS exposed to PWR primary environment indicates that almost 80% of the Operational Experience (OPEX) occurred in low flow or stagnant zones in dead leg situations where the primary water chemistry was probably contaminated by anionic impurities especially trapped oxygen bubbles in the case of Canopy Seals. Further, the radiolysis in closed volumes such as blind-end holes with limited or no ingress of conditioned hydrogenated PWR primary coolant may lead to increased REDOX-potentials. Oxygenated PWR primary water transients could also be a potentially important factor in the in-service failures that have occurred in cold worked stainless steels in PWRs. Indeed, oxygen ingress can occur under a variety of circumstances such as load following if aerated feed water is injected while deborating the primary coolant.

A dedicated program based on SSRT tests was designed to evaluate the PWSCC susceptibility at 345°C of type 316L austenitic stainless steel. Three different environments were chosen; two oxidizing PWR primary environment with 10 ppm and 0.1 ppm of dissolved oxygen and a reducing nominal PWR primary environment with a hydrogen partial pressure of 0.3 bar.

Keywords: 316L stainless steels, PWSCC, SSRT

1. INTRODUCTION

The excellent resistance of austenitic stainless steels to PWSCC is well admitted by the international community, however in specific material and environmental conditions, degradation can occur. This latter consideration is particularly apparent from the recent EPRI review [1]. Indeed, cold working driven by its strain history plays an important role on the SCC phenomenon. However, this material aspect cannot be considered on its own without taking into account the synergy with abnormal water chemistries.

In this context, oxygenated PWR primary water transients could be a potentially important factor regarding the in-service failures that have occurred in cold worked stainless steels in PWRs. Indeed, oxygen ingress can occur under a variety of circumstances [2-4]. For example, during load following, the initial step requires the addition of boric acid to moderate the reaction. In order to increase the power to 100%, the primary water is diluted with the reactor make-up water via the Chemistry and Volume Control System (CVCS) charging line. Such tanks can be partially or fully aerated. Oxygen added with the reactor make-up water has been suggested to be the cause of chemistry anomalies in some load following


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measured PWRs. The actual concentration of oxygen in the Reactor Coolant System (RCS) depends mainly on the CVCS and RCS flow rates. Some of this injected partially or fully aerated water can also go directly to the pressurizer spray nozzles and primary pumps. Such components, partly made of CW stainless steels may be at risk for SCC, therefore some efforts are needed.

2. MATERIAL OF INTEREST

The austenitic alloy tested in this work is an AISI 316L (AZ2 CND 18-12) 21-mm thick plate heat T8289 manufactured by hot rolling followed by a water quench. Chemical composition and mechanical properties of the material are respectively presented in Table 1 and Table 2 and fit with the French RCC-M 3307 requirements. The martensite start temperature Ms is -224°C (see Table 3), while the temperature Md30 at which 50% of martensite will be formed after a true straining of 30% is -7.2°C, as predicted by Angel [5]. Consequently, the probability to obtain a strain induced martensite is rather low. The amount of residual ferrite (3.3%) was evaluated using an empiric law developed for austeno-ferritic stainless steels and where xi is the fraction number of element i [6].

With

18.508.03.820

6.17065.0

MnNCNi

SiMoCr

xxxx

xxxR

Stacking fault energy (SFE) of stainless steels is related to the chemical composition of the material and can be evaluated using the following empiric law [1]:

MnSiNCrCNi xxxxxxmmJSFE 2.113779.041027.25).( 2

A SFE of 30.4 mJ•m-1 was calculated. This value is within the typical range measured on stainless steels.

The microstructure did not contain any carbide precipitates in the matrix or along the grain boundaries (Figure 1). The microstructure was analyzed by EBSD (Figure 2 and Figure 3). Results supported by pole figures presented in Figure 4 indicate a marked texture with a maximum texture index of 4.5.

3. EXPERIMENTAL PROCEDURE

3.1. Specimens

Two different types of specimens were used in this program, V-humped and flat peened specimens.

V-humped specimens were made from flat strips machined by EDM in the TS orientation relative to the material processing conditions (SCC convention) of the T8289 material. Flat strips were then manually and gradually polished using silicon carbides papers up to #1200 to obtain a thickness of 1.4 mm. The hump, as shown in Figure 5, was introduced in the middle of the gauge length using a dedicated device. No mechanical cracks were detected on the V tip using optical microscopy.

A HV0.1 hardness map was performed on a cross section of a humped specimen focusing on the humped area in order to evaluate the cold work induced by the forming process. Results indicate that the hardness

39.396.5²8.21(%) RR


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in undeformed areas is close to 180 HV0.1. The maximum hardness of 350-380 HV0.1 was symmetrically located close to the V tip, whereas the tip exhibited a lower hardness of 260 HV0.1. The resulting strain gradients that might be developed during the test at the V tip could promote the PWSCC initiation.

The V-humped specimen was slow strain rate tensile tested as described in the next section. In the frame of these SSRT tests, since in static conditions RUB specimens have demonstrated their susceptibility [7], it was decided to test an intermediate case having the hardness properties of RUB specimen but tested in SSRT conditions, originally supposed to be more severe.

Flat tensile peened specimens (50 mm long and 3mm width were prepared in a similar way to the humped specimen up to the final polishing step reaching a thickness of 1.4 mm with a silicon carbide paper #1200. One surface of the specimen was then shot peened for 2 minutes with a pressure of 2 bar on the whole surface with 200 μm diameter balls made of corundum at a distance of 20 cm. The shot peened flat strip was then tensioned up to a total strain of 21.7% corresponding to the average through thickness strain of a RUB presented in simulated by FEM.

HV0.025 profiles were collected on a cross section of a flat peened specimen in order to evaluate the cold work level induced by the combined effect of the peening and the tension. The hardness data are presented in Figure 7 with data collected at the apex of a RUB presented in [7]. These results indicate very similar trends with a surface hardness higher than 400 HV0.025 progressively reaching plateau hardness within the 240- 260 HV0.025 range at a depth close to 100 µm. This change of slope corresponds to the transition between surface cold work induced by the peening followed by the tensile test, and bulk cold work only due to the tensile test.

3.2. Test apparatus

The following test matrix (Table 4) summarizes the 4 different conditions tested within this program.

Before exposure, specimens were ultrasonically rinsed in ethanol and then in distilled water. Tests were carried out in stainless steels autoclaves. Specimens were insulated from the autoclave by oxidized Zircaloy to avoid galvanic coupling.

The two first SSRT tests using both the flat peened and a V-humped specimen were conducted in a solution of pure water containing 2000 mg/kg of boron as boric acid and 0.1 mg/kg of lithium as lithium hydroxide, conservatively simulating PWR primary shutdown conditions at a temperature of 345°C, in an autoclave coupled with a refreshing loop. In this condition, the amount of oxygen was controlled and kept at 10 mg/kg H2O.

Following the results of these two first tests, an additional V-humped specimens were SSRT tested in a solution of pure water containing 1200 mg/kg of boron as boric acid and 2 mg/kg of lithium as lithium hydroxide at 345°C in autoclaves coupled with refreshing loops. These two additional second tests were respectively conducted with an amount of oxygen kept at 0.1 mg/kg H2O.

A third set of tests were conducted in a deaerated environment consisting in a solution of pure water containing 1200 mg/kg of boron as boric acid and 2 mg/kg of lithium as lithium hydroxide at 345°C in autoclave with a hydrogen partial pressure of 0.3 bar (65 cc/kg of H20).

The water chemistry was analyzed and validated before and after each test (B, Li, Cl-, SO42-,F-). In all the

tests, the crosshead speed was fixed, and for the flat tensile specimens taking into account the load train compliance, this produced a strain rate of 2x10-7 s-1. The local strain rate at the V tip of the humped specimens was calculated by FEM and was found to range between 3x10-7 s-1 and 5x10-7 s-1 depending on the V opening state. Consequently both strain rates are within the same order of magnitude.


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4. RESULTS AND DISCUSSION

Results of the four SSRT tests are presented in Table 5 highlighting the SCC characteristics in particular both the maximum crack depth and the crack path.

SEM observations performed on specimen EAC6-1 tested in shutdown PWR primary environment with 10 ppm of dissolved oxygen clearly indicate the presence of intergranular (IG) fracture covering the whole fracture surface. A similar characterization performed on the specimen EAC6-3 (Figure 9) tested at 0.1 ppm of dissolved oxygen had a transgranular (TG) fracture surface (Figure 10) before a final mechanical ductile fracture of the specimen. The presence of striations on the fracture surface can also be seen (Figure 11). Such striations were identified in the literature as crack arrest points [8].

After a test duration of 2000 h, the specimens EAC6-2 (nominal PWR primary water conditions) was removed from the autoclave. SEM examinations of the inner surface of the V presented in Figure 12, indicate a high density of shallow cracks. Further examinations of the outer surface of the V indicate the presence of short SCC cracks on both shoulders. Detailed surface analysis indicates that almost 90% of the cracks are TGSCC and 10% IGSCC. TGSCC initiated from the emergence of shear bands at the surface preferentially oriented perpendicularly to the loading axis (Figure 13) whereas IGSCC was mainly observed between grains presenting visible strain incompatibilities and also preferentially orientated perpendicularly to the loading axis (Figure 14). Statistical evaluation of the crack depth as a function of the distance to the V tip performed on a cross section is presented in Figure 15. Results indicate that the maximum crack depth was reached at 200 µm from the V-tip corresponding to the area where the highest initial hardness gradient was measured (Figure 16). The highest density of the deepest shallow cracks was also observed in the area presenting an initial hardness higher than 280 HV0.1. An important density of shallow cracks were also noticed at the outer surface of the V were the initial hardness was higher than 240HV0.1.

In these particular tests, it seems that the SCC initiation susceptibility was more governed by the initial strain gradient induced by the forming process rather than the level of strain. This means that interfaces between “hard zones” and “soft zones” corresponds to the preferential crack initiation site (Figure 16).

The change in cracking mode took place between 10 and 0.1 ppm oxygen. Additional tests within this dissolved oxygen (DO) range will be needed to understand whether the transition is gradual or whether it occurs rapidly at a particular DO level. However, it is clear that for these severe loading conditions including humping and dynamic straining far away from field reality, it seems that only small amount of oxygen is required to drastically change the PWSCC susceptibility of cold work stainless steels. Therefore, the risk of primary water chemistry transients containing significant amounts of dissolved oxygen has to be carefully studied, especially in the case of high flow rate can drastically increase the electrochemical potential [9-10].

The flat tensile peened specimen referenced 316-SPH-T1b tested in shutdown PWR primary environment with 10 ppm of dissolved oxygen failed after 9% of elongation. SEM examinations of the fracture surface presented in Figure 17 indicate a fully ductile failure of the specimen with the noticeable presence of TGSCC initiated from the peened surface. Furthermore, SEM observations of a cross section indicate the presence of numerous blunted TGSCC cracks (Figure 18) with an average depth of close 100 µm corresponding to the hardness transition presented in Figure 7. Locally, some TGSCC cracks were associated with lateral intergranular oxide penetrations (Figure 19) that might correspond to TG/IG transition precursors. Despite this observation, it seems that in this particular test, the cracks did not propagate after the first step of initiation.


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In the case of the peened RUB specimens presented in Figure 20 [7], a similar TG crack initiation trend was observed within the first 150 µm, but contrary to the peened tensile specimen, the crack did propagate intergranularly through the whole thickness for a very similar initial cold work and passive load.

The difference in PWSCC susceptibility between the flat tensile peened specimen and the RUB could be related to an inadequate loading strain rate in the case of the SSRT test. However, it is worth noting that the V humped specimen is highly susceptible to SCC in a very similar strain rate range.

Stresses analyses performed by XRD and FEM showed that RUB and flat shot peened specimens developed stresses (800-1000 MPa) higher than V humped specimen at the crack initiation site location (500-700 MPa), suggesting that the strain gradient was a major factor in the SCC susceptibility. As a consequence, among the usual discussed criteria related to SCC occurrence, such as hardness and stresses, strain gradient should not be neglected.

5. CONCLUSION

The analysis of the recent field experience on PWSCC of cold worked stainless steels highlights the important role of off-normal chemistry and led the development of experimental programs to develop maintenance strategies. In response to this, AREVA launched an internal R&D program to better understand SCC in such considerations.

The detrimental effect of oxidizing environment was studied in several different oxidizing environments at 345°C. The following conclusion can be drawn:

In s nominal PWR primary water environment at 345°C, the application of dynamic loading on V-humped specimens led to TG shallow cracking. In contrast, 10 ppm of DO led to a complete IG specimen failure. An additional test using an intermediate content of 0.1 ppm of DO led to a TG response. It seems that only 0.1 ppm DO has a very important effect on the PWSCC susceptibility of cold worked 316L SS at 345°C.

The large difference in SCC response between the flat peened tensile specimen and the V-humped specimens suggests that the latter specimen type is highly appropriate for assessing PWSCC susceptibility of stainless steels. It seems that the strain history developed during forming and loading steps, especially strain gradients, might play a key role in SCC occurrence and has to be carefully taken into account.

References

[1] G. O. Ilevbare, F. Cattant, N. K. Peat, “SCC of Stainless Steels under PWR service conditions”, Proceedings of International Symposium Fontevraud 7, Avignon, France, 26-30 September 2010.

[2] P. Combrade, F. Vaillant “Corrosions des alliages de Nickel et des aciers inoxydables en milieux REP polluées ou confinés » Dossier Technique de L’ingénieur, bn3756, 10 January 2014.

[3] P. Millet « PWR primary water chemistry guidelines » Volume 1, Revision 4, EPRI, TR-105714-V1R4, 1999.

[4] R. Roofthooft, et al., “Primary Circuit Chemistry of Western PWR and VVER Power Plants” 02004Ren9653, p. 69, Paris, France, 1996.

[5] F.B. PICKERING, "Physical metallurgy and the design of steels", Applied Science Publishers, London, 1978.


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[6] J.P. MASSOUD, M. BETHMONT, J. CHAMPREDONDE, ”Long term aging of cast duplex stainless steel between 300 and 400°C Relationship between toughness properties and metallurgical parameters, Duplex Stainless Steels'91” , Vol. 1,p. 93, 1991,

[7] G.Ilevbare, “Experimental Study of Stress Corrosion Cracking Initiation in Austenitic Stainless Steels in Off-Normal Chemistry PWR Primary Water Environments” EPRI report 3002000552, 2013.

[8] T. Couvant, L. Legras, F. Vaillant, J.M. Boursier, Y. Rouillon, ‘Effect of strain hardening on stress corrosion cracking of AISI 304L stainless steel in PWR primary environment at 360°C’, 12th International Conference on environmental degradation of materials in nuclear systems-water reactors, Snowbird (Utah), 2005.

[9] D. MacDonald, “Viability of Hydrogen Water Chemistry for Protecting In-vessel Components of BWR”, Corrosion, 1992

[10] P. Ford, P. Scott, P. Combrade “Environmentally Assisted Degradation of Stainless Steels in LWRs” ANT international, 2010.


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Tables

Table 1: Chemical composition of T8289

Material Heat C (%) Si

(%)

S

(%)

Mn (%)

Ni (%)

Cr (%)

Fe (%)

Mo (%)

P

(%)

B (%)

Cu (%)

N (%)

316L T8289 0.023 0.488 0.008 1.73 12.31 17.31 2.45 0.025 0.132 0.07

RCC-M M3307 <0.035 <1 <0.015 <2 11.5

- 12.5

17 -

18 -

2.25 -

2.5 <0.03 - <1.00 -

Table 2: Mechanical properties of T8289

Material Heat Temperature (°C) Ys (MPa) UTS (MPa) A %

316L T8289 20 270 575 62

350 NA NA NA

RCC-M M3307 20 >220 >520 >40

350 >130 >400 -

Table 3: Metallurgical characterizations

Material Heat Ms (°C) Md30 (°C) Grain size (µm) Ferrite % SFE (mJ/m)

316L T8289 -224 -7.2 125 3.3 30.4

Table 4: Test matrix

ID Heat Specimen Temperature PWR primary

environment

Hydrogen

partial

pressure

Dissolved

oxygen

content

316-SPH-T1b

T8289

Flat Peened

345°C

2000 ppm B 0.1 ppm Li -

10 ppm EAC6-1

V-humped EAC6-3 1200 ppm B 2ppm Li

0,1 ppm

EAC6-2 0,3 b <5 ppb


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Table 5: Results obtained on SSRT tests

ID Specimen Temperature DO Duration

(h)

Strain rate (s-1)

a IGSCC*

(µm)

a TGSCC**

(µm)

SCC % ***

Crack feature

IGSCC %

TGSCC%

316-SPH-T1b

Flat Peened

345°C

10 ppm125 2x10-7 90 6 5 95

EAC6-1

V-humped

930 3x10-7

to

5x10-7

1400 100 100 -

EAC6-3 0,1ppm 1150 500 60 0 100

EAC6-2 <5ppb 2000 60 5 10 90 a IGSCC* represents the maximum Intergranular Stress Corrosion Crack depth observed a TGSCC** represents the maximum Transgranular Stress Corrosion Crack depth observed SCC%*** represents the ratio between the maximum SCC crack depth and the specimen thickness


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Figures

Figure 1: SEM micrograph, T8289, TD

Figure 2: Inverse pole figure, T8289, RD, plane TL

Figure 3: Inverse pole figure, T8289, TD plane TL

Figure 4: Pole figure, T8289, RD


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Figure 5: V-humped specimen Figure 6: HV0.1 map of a V-humped specimen

Figure 7: HV0.025 profiles of a flat peened specimen and a peened RUB

100

150

200

250

300

350

400

450

500

0 0.05 0.1 0.15 0.2 0.25 0.3 0.35 0.4 0.45 0.5

HV

0.02

5

Distance from the peened surface (mm)

Flat peened and 20% tensioned specimen

Peened RUB


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Figure 9: fracture surface overview, EAC6-3, 1150 h, 345°C, 0.1 ppm DO

Figure 10: TGSCC fracture surface, EAC6-3, 1150 h, 345°C, 0.1 ppm DO

Figure 8: IGSCC fracture surface, EAC6-1, 930 h, 345°C, 10 ppm DO

Crack propagation

Crack propagation

Crack InitiationCrack Initiation

Crack propagation


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Figure 11: Striation observed on TGSCC fracture surface, EAC6-3, 1150 h, 345°C, 0.1

ppm DO

Figure 12: V Inner surface overview, EAC6-2, 2000 h, 345°C, 0.3 bar pH2

Figure 13: TGSCC/ IGSCC crack initiation, EAC6-2, 2000 h, 345°C, 0.3 bar pH2

Figure 14: strain incompatibilities at grain boundary leading to IGSCC initiation, EAC6-2,

2000 h, 345°C, 0.3 bar pH2

Plane trace

IGSCC

TGSCC

Crack propagation


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Figure 15: Cross section of EAC6-2, 1500h, 345°C, 0.3 bar pH2

Figure 16: V-humped specimen, HV0.1

0

10

20

30

40

50

60

0 1 2 3 4 5 6 7 8

Distance from the V tip (mm)

Cra

ck d

epth

(µm

)V Inner Surface

V Outer Surface

Deepest crack


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Figure 17: Fracture surface overview, 316SPH-SP-T1b, 9% strain, 345°C, 10 ppm DO

Figure 18: Cross section of 316SPH-SP-T1b, 9% strain, 345°C, 10 ppm DO, blunted TGSCC cracks


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Figure 19: Cross section of a peened RUB,1500 h, 345°C, 10 ppm DO, IG preferential

oxidation

Figure 20: Cross section of a peened RUB,1500 h, 345°C, 10 ppm

DO, TG/IG transition

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STRESS CORROSION CRACKING OF STAINLESS STEELS TESTED …envdeg2015.org/final-proceedings/ENVDEG/papers/ENVDEG126.pdf · (1) AREVA NP Technical Center, 30 Boulevard de l'Industrie