Hydrogen embrittlement in the duplex stainless steel Z2CND2205 hydrogen-charged at 200°C

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Materials Science and Engineering A224 (1997) 116-124

Hydrogen embrittlement in the duplex stainless steel Z2CND2205 hydrogen-charged at 200°C F. Iacoviello a,,, M. H a b a s h i b, M. C a v a l l i n i c Universit~ di Cassino, Dipartimemo Ingegneria Industriale, via G. di Biasio 43, 03043 Cassino (FR), Italy u CNRS-Ecole Centrale Paris, Laboratoire CFH, Grande Voie des Vigncs, Chdtenay-Malabry 92295, France ° Umc'ersir~ di Roma La Sapienza, Dlparttmento ICMMPM, via Eudossiana, Rome, Ball,

Received 21 February 1996; revised 23 September 1996

Abstract

Physical and mechanical properties of austeno-ferritic stainless steels depend on the microstructure and phase transformations: many intermetallic phases, carbides and nitrides precipitate at different tempering temperatures. Hydrogen behaviour in steels is affected by the morphoIogy and by the presence of precipitates, both for its diffusionai behaviour and for the importance of trapping phenomena. In this paper, the hydrogenation of a 22Cr-SNi duplex stainless steel has been achieved at 200°C in molten salts bath in potentiostatic conditions, and hydrogen embrittlement has been characterised using low strain rate tensile tests. The influence on hydrogen embrittlement of different intermetallic phases, carbides and nitrides has been considered via tempering heat treatment with tempering temperatures between 200 and 1050°C. The possibility of the recovery of the mechanicaI properties of charged steel outgassing at room temperature has been considered, investigating also the influence of different intermetallic phases, carbides and nitrides. Moreover, a new thermal technique based on high-temperature outgassing tests has been performed in order to calculate the hydrogen coefficient of diffusion. © 1997 Elsevier Science S.A. Keywords: Carbides; Hydrogen embrittlement" Intermetallic phases; Nitrides; Stainless steel

1. I n t r o d u c t i o n

Austeno-ferritic duplex stainless steels are widely used in oil and nuclear industries due to their reasonable costs, good mechanical and thermal properties, and resistance to stress corrosion and pitting. However, their mechanical properties are usually reduced by the hydrogen presence, which could increase the probability of failure in service [1-5]. In the austeno-ferritic stainless steels, after tempering between 350 and 500°C, an age hardening of the ferrite is observed, the result of which is an embrittlement at 475°C. This embrittlement is due to two different mechanisms: (a) The first mechanism is a spinodal decomposition of the ~-ferrite in two phases, an a'-Cr-rich phase and c~-Fe-rich phase, in the temperature range 350475°C. This is a non-diffusional mechanism [6-8].

* Corresponding author. 0921-5093/97/$17.00 © 1997 Elsevier Science S.A. All rights reserved. PII S0921-5093(96)10545-I

(b) The second mechanism is nucleation and growing of G-phase (about 500°C). Cr, Mo and Cu increase the sensitivity to the hardening at 475°C. Ni promotes the distribution of the Cr, Mo and N in the c~-phase and also accelerates the spinodal decomposition. The G-phase, rich in Ni, Si, Mn and AI, forms at 350°C for an ageing time between 1000 and 30 000 h [9]. During the isothermal heat treatment between 600 and 1050°C, many phases develop [10-12]: ( a ) y - p h a s e (600-900°C), due to the decomposition of c%phase: (5--+ ~ ( 7 0 0 - 9 0 0 ° C ) + y, or due to a 'martensitic' transformation mechanism; (b) ~-phase (700-900°C), a tetragonal structure, compact and complex, hard and brittle: this phase precipitates at the triple joints of the grain boundaries or at the austenite-ferrite interfaces; (c) other phases: R (550-700°C), X (700-9000C) and 7z [10-12]; (d) precipitation of carbides such as M23C6, MvC 3 and Cr2N in the temperature range 600-1100°C [10,11]. The purpose of this work is to investigate the effects on the mechanical properties of the interactions be-

F. Mcoviello et al./Materials Science and Dzgineering A224 (I997) 1 I6-124

II7

Table 1 Duplex stainless steel Z2CND2205 chemical composition Element

C

Si

Mn

S

P

Ni

Cr

Mo

Cu

Co

N2

02 (ppm)

Wt.%

0.025

0.385

1.428

0.01i

0.028

5.639

22.78

2.491

0.148

0.160

0.129

49/54

tween the hydrogen atoms and the different phases that can form after 3 h tempering at between 200 and 1050°C. Taking into account that the hydrogen diffusion coefficient in the duplex stainless steels at T = 200C is very low (DH = 10 - to cm 2 s - 1) [13], the duplex stainless steel Z2CND2205 has been charged in a salts bath at 200°C (DH = 1 0 - 7 c m 2 s - 1 ) [13]. In this investigation, we have improved this finding by applying a new thermal technique. In order to quantify the hydrogen embrittlement of this steel, low strain rate tensile tests ( ~ 10 .6 s -I) have been performed for different conditions of tempering temperature, after cathodic charging. The outgassing of hydrogen-charged samples has been carried out under v a c u u m ( 1 0 - 3 Pa) at 600°C, in order to measure the desorbed quantity of hydrogen, Qr~, after each specific heat treatment. Samples outgassed at 600°C have also been outgassed to fusion using another experimental device: residual hydrogen quantities of 3-5 ppm have been obtained. This residual hydrogen quantity is not dependent on the tempering temperature. Outgassing tests have also been carried out at a constant temperature in the range 150-6000C, to measure the hydrogen diffusion coefficient by performing the thermal technique. This method has been compared with the permeation one. Our method has been applied firstly on austenitic AISI A286 steel and then on the Z2CND2205 steel. Moreover, the microstructure and the crack surfaces have been observed using both an optical microscope and a scanning electronic microscope (SEM).

are 60 and 40%, respectively, and have been measured using the X-ray diffraction technique. Tempering heat treatments have been carried out for 3 h between 200 and 1050°C, at almost every 50°C, in argon atmosphere, followed by water quenching. 2.1. Cathodic charging conditions at 200°C

The cathodic charging at 200°C has been carried out by applying - t . 5 V per Ag (potentiostatic polarisation) in a molten salts bath (sulphates and bisulphates of Na and K), for different time periods: 0.5, 1, 3 and 20 h. The presence of hydrogen is assured by the hydrolysis of the water dropping in the molten salts bath. Fig. l(a) and (b) show that, for a steel tempered at 200°C, the ductility at failure decreases with increas!!

1100 / ~

9OO

(Y

(lVlPa)

J l

Not charged

700 20

3

1

0.5

5O0

'

300

0

' 0.02

'

'11 ' 0.06 0.1

0.04

' 0.3

0.2

E

0.6 1

3

t c (h)

100

20

100

2. Material

The chemical composition and tensile properties of the investigated Z2CND2205 steel are shown in Tables 1 and 2. This steel corresponds to the commercial steel 'Uranus 45N' and is supplied as cold rolled to 1 mm thick. The percentage contents of ferrite and austenite Table 2 As-roiled duplex stainless steel Z2CND2205 tensile properties YS (MPa)

UTS (MPa)

~m (%)

725__+25

885_+ 10

18.6+ 1.4

0

0 O

100

200

300

i

Fig. I. (a) Tensile curves for different charging times (h) in salts bath at 200°C; (b) QH and F% variation as a function of ~/':t~.

118

F. [acoviello et al./ Materials Science and Engineering A224 (1997) 116-124

ing charging time. Nevertheless, the resistance .Rm is not too sensitive to the charging time (Fig. l(a)). Fig. l(b) represents the variation of QH (hydrogen quantity desorbed at 600°C) and }7% (i.e., relative ductility loss with and without hydrogen). The two parameters depend on x/to (where G is the charging time). It is evident from Fig. l(b) that there is a linear relationship between Q~ and x/to and that t7% becomes constant for to >~ 3 h. Hydrogen charging is influenced by both the hydrogen diffusional behaviour in austenite and ferrite grains, and the hydrogen trapping. Dislocations and grain boundaries are main trapping sites for the duplex stainless steel in as-rolled conditions. These trapping sites are characterised by a reversible behaviour, with a continuous releasing of hydrogen during degassing experiments, while irreversible ones do so after a critical temperature [14,15]. Considering the relationship t = ea/ 4D~ (i.e., e is the specimen thickness, and DH is the hydrogen diffusion coefficient) at 200°C, the time necessary to saturate the interstitial sites of a 1 mm thick duplex stainless steel specimen (i.e., G) is between 3.5 and 7 h. Charging times higher than t~ (such as 20 h) correspond to an increase of the saturation of the existing trapping sites, together with the formation of new trapping sites (such as microvoids, characterised by the same interaction energy of dislocations and grain boundaries) and the saturation of interstitial sites in ferrite and austenite. At 600°C, it is possible to outgas both interstitial sites and trapping sites, respectively saturated and not saturated, after a charging time of t~. On the basis of these results, we have chosen 3 h as the charging time for all the heat treatment conditions carried out in this study. Before describing the experimental results concerning the effects on the mechanical properties of the change in microstructure in the absence and presence of hydrogen, we gve below a brief description of the theoretical basis of the thermal technique used here to measure the hydrogen coefficient of diffusion in the range 150-600°C.

2.2. Hydrogen difJi~sion coefficient (DH) measurements

The diffusion coefficient DH in a metal depends on its crystalline system (b.c.c., f.c.c., h.c., etc.), on its microstructure (ferrite, martensite, bainite, etc.), on the presence of traps such as dislocations, dispersed phases, grains boundaries, interfaces between phases, etc., and finally, on the test temperature. B.c.c. (a) and f.c.c. (?,) structures have a completely different hydrogen diffusional behaviour, with differences in the diffusion coefficient values (for example, at room temperature, DH~ is approximately equal to 3-5 times DH:.), and also in activation energies (EH~ z~FHT; EH is the energy barrier that must be overcome by thermal fluctuations in the system). ----

The relationship between the hydrogen diffusion coefficient (DH) and the test temperature (T) can be described by the Arrhenius law:

l

(1)

where Do is a constant and EH is the activation energy of diffusion, the value of which depends on the parameters described above. Hydrogen diffusion coefficient values are measured using an outgassing method performed at several constant temperatures under vacuum (10 -3 Pa). This method allows the hydrogen diffusion coefficient value to be calculated by measuring the quantity of the outgassed hydrogen at a constant temperature as a function of the outgassing time on a constant thickness of a plane sheet, with the assumption that this coefficient is constant during the test. The value obtained is only apparent, taking into account diffusion in both austenitic and ferritic phases. This method has been used to determine, for example, the diffusion coefficient of direct dyes in cellulose sheet [16] and oxygen in muscle [17]. Defining QH(t)/QH(t= co) as the normalised outgassed hydrogen quantity (the ratio between the hydrogen quantity outgassed after a time t and the hydrogen quantity outgassed after an infinite time), and considering that the hydrogen diffusion coefficient is hydrogenconcentration-independent, the appropriate solution of the diffusion equation may be written as following [18]: QH(t)

8

Q (t = co)= 1

~"

1

1)

x exp{ -- DH(2ml+e 2 )27t2(}

(2)

where e is the specimen thickness and D , is the hydrogen diffusion coefficient. Considering the time t0.5, corresponding to a normalised outgassed hydrogen ratio QH(t)/QH(t = co)= 0.5, the previous equation gives:

with an error of about 0.001%. This relationship can be simplified by: D

-

0.049e a -

-

to.5

(3)

Eq. (3) then allows the hydrogen diffusion coefficient to be calculated, knowing the specimen thickness (in this work, e = 0.1 cm) and the outgassing curve QH -- t. This method is related to the outgassing rate: if the outgassing rate is high (high diffusion coefficient), the

F. Iacoviello et a l . / Materials Sczence and Engineer#zg A224 (i997) i I 6 - I 2 4

[°C]

T 600

10 "9

2o0

400

20 []

El

N

10-I0

~

authors ref. [i9] ref. [20]

10 -1I

\"

<,,, ,,<,= \

10 -I2

\ \

\

10 -I3

\ \

\ \

10 -I4

, 1.0

I 1.5

,

I

,

2.0

1000/T

I

,

2.5

I

,

3.0

3.5

[K 1]

Fig. 2. H y d r o g e n diffusion coefficients (DH) vs. temperature: A I S I A286 steel.

variation of the outgassed hydrogen is difficult to measure. With our experimental equipment, we have been able to measure hydrogen diffusion coefficients up to 7 x 10-*° m 2 s-1. On the other hand, an excessively slow outgassing rate is related to low diffusion coefficients, which implies that the test period is too long. It is worth noting that this method has been used to calculate the hydrogen diffusion coefficients, not only on the austenite AISI A286 and duplex steel Z2CND2205, but also on nickel alloys such as 600, 690 and 800, and on AISI 321 stainless steel (these results will be published elsewhere).

119

Fig. 2 shows that there is a good agreement between the two techniques, especially with the results of Quick and Johnson [20], in which the activation energy EH is equal to 49 kJ mol - t while our results give 42 kJ molThe results obtained on the duplex stainless steel Z2CND2205 (Fig. 3) have also been compared with those obtained using the permeation technique [1.19,13], on AISI 301 with ~ ' = 0.6 in the range 170320°C [19] (where a' is the volume fraction of residual martensite in the austenitic stainless steel), duplex steel 2205 at 20-350°C [1], duplex stainless steel 2206 in the range 22-80°C [13] and ferritic steel A L 29-4-2 in the range 100-350°C [19] (Fig. 3). The results obtained on AISI A286 steel are also plotted. The variation of DH as function of lIT on the duplex stainless steel studied in this work is also in good agreement with those of the other authors. Fig. 3 also shows that the value of the hydrogen diffusion coefficient of the steels containing two phases, such as ),-c~ or ?-cd, is situated between those of D H. . . . , and DHr. The D~ . . . . , is always higher than DH~(~,+~) and DHy. Thanks to its microstructure, the hydrogen diffusion coefficient of duplex stainless steel Z2CND2205 follows a simple series model, which allows the hydrogen diffusion coefficient of a and ?,, and their volume fraction, to be correlated with the hydrogen diffusion coefficient of the duplex stainless steel. This model is a simplification of the one presented by Bernabai and Torella [21] and it can be written as follows:

1

f~

DR

fr

-- DH-""~ 4- DH'-"" ~

(4) T

600

3. Experimental results and discussion

16s

Firstly, we discuss the results of the hydrogen diffusion coefficient measurements on AISI A286 and duplex Z2CND2205 steels, and secondly, we give the results of the mechanical properties of the duplex stainless steel Z2CND2205 with different microstructures in the absence and presence of the internal hydrogen.

~(59

3.1. Hydrogen diffusion coefficient measurenwnts The hydrogen diffusion coefficient of an annealed (1050°C h - 1 ) AISI A286 stainless steel has been measured in the temperature range 200-600°C. This steel has been chosen for its stable austenitic structure, which allows this method to be improved by comparing our results with those obtained by other authors [19,20] on the same steel in annealed conditions, and by applying the permeation technique at different temperature ranges (100-350°C [19], 200-500°C [20]).

400

[°C]

200

20

8

\ N N

""

1(51C

% % \ \

"v

\

An

1611 16t:

\ \

AISI A286 [authors] a

Z2CND2205 [authors] AL 29-4-2 ferritic steel [19]

1(5 I2

zx

1(5i 4

V,

• 1.0

,,

AISI301 withe.V=60% [19] Duplex 2206 with o: =56% [11] IDup{ex 2205 [1],

1.5

2.0

1000/T

,

2.5

,

,

3.0

,

~, 3.5

[K 1]

Fig. 3. H y d r o g e n diffusion coefficients vs. temperature: duplex steel Z2CND2205.

120

F. Iacov'wllo et al./ Materials Science and Eng#wering A224 (i997) i 16-I24

1100

1100

1000

1000

500

500

f-

200

600

I

200 700

800 900 Re et Rm (MPa)

1000

1100

0

I

!

10

20

25

Fig. 4. R~, R m et % ~ariation with the tempering temperature: uncharged condition. The differences between the experimental values and the model are included in the experimental scatter. 3.2. Mechanical properties in the absence and presence of internal hydrogen 3.2.i. Mechanical properties withozlt hydrogen Fig. 4(a) and (b) shows the variation of Re, Rm and er (average values) as a function of the tempering temperature, without hydrogen. These figures show two maxima both in Re and Rm at 475 and 800°C, corresponding, respectively, to the spinodaI decomposition ~--+ c~' and the precipitation of a-phase (Fig. 5). Fracture surface morphologies change completely according to the tempering temperature, from a completely ductile to a totally brittle one (Figs. 6 and 7). A minimum in R~ and _Rm at 750°C, probably due to the formation of the ), and M~3C 6 (Fig. 8) phases [10,11], has been also observed.

Fig. 5. Formanon of o--phaseand secondaryaustenite (after ref. [10]).

Superimposing the c u r v e .R m - T to the T T T diagram obtained by Van Nassau et al. [12] (Fig. 9), it is possible to verify that the variation of R m corresponds effectively to the precipitation and formation of several phases durhag the tempering. 3.2.2. Mechanical properties with hydrogen The variation of the yield strength Re with the tempering temperature is the same as that measured without hydrogen (Fig. 4(a)). Moreover, the variation of Rm with the temperature in the presence of hydrogen is the same of that obtained without hydrogen, with a mean shift about 150 MPa. In Fig. 9, the T T T diagram of the duplex stainless steel Z2CND2205 [12] has two C curves: the first is between 300 and 500°C, and the second is between 500 and 1000°C. We can conclude, then, that the mechanical properties are very sensitive to microstructure variation. For the second C curve, the different trapping sites of the hydrogen atoms due to the precipitation of these phases (cV, 7, a, M~Cy) are observed not only in the mechani-

Fig. 6. Tensile fracture surface for a 3 h tempering heat treatment at 1050°C (uncharged conditions): completelyductile morphology.

12I

F. lacoviello et al./ Matertals Science and Engineering A224 (I997) II6-I24

1000

900 800 700 600

5" % [-.

500 Rm(MPa)

400

600

700 i

30O

800 1

900~ I ~

@

1000 I

1t00 I

!

'

'

1

~

1'o

'

1200 I

1~o

t r (h)

Fig. 7. Tensile fracture surface for a 3 h tempering heat treatment at 800°C (uncharged conditions): completelyfragile morphology. cal properties, but are also reflected in the variation of QH as a function of 2 (2 = Rwx4/RmAIP,) and the tempering temperature ( 2 - T and 2-QH) (Fig. 10). The choice of f~m as a parameter eharacterising the hydrogen embrittlement has been made because of its sensitivity to both hydrogen embrittlement and to the precipitation of different intermetallic phases, carbides and nitrides. Moreover, tempering temperatures corresponding to the second C curve show a reduction of area and an elongation too affected by the precipitation of different intermetallic phases, carbide and nitrides, and it could be very difficult to distinguish the hydrogen embrittlement effects. Fig. 11 shows the variation of 2 and QH for each tempering temperature (second C curve). The relationship between 2 and QH is linear and the ,;-values decrease with decreasing QH, the lower 2-values corresponding to higher embrittlement. For the tempering temperature of 800°C, where the ~r-phase volume fraction is optimal [10], 2 = 70% and the QH-value is low (17.5 ppm). On the other hand, for 1050°C, where the phases are practically dissolved, 2 = 90%, in spite of a large Qn-value of about 40 ppm.

Fig. 9. Superposition of the Rm-T curve wtth TTT dmgram [12]: duplex stainless steel Z2CND2205. It is possible now to evaluate the sensitivity of the microstructure of the duplex steel to hydrogen for each value of 2 between 70 and 90%. In Figs. 12 and 13, it is possible to note the difference of the tensile fracture surface corresponding to the two minima in the T T T diagram for charged conditions. Fig. 14 shows the variation of 2 and QH for the first C curve. This behaviour is explained as follows: for temperatures in the range 200-550°C, 2 is high (low embrittlement). However, the experimental results are divided into two groups: for T/> 350°C, with the exception of T = 400°C, the QH-value is low with respect to that obtained for T = 200 and 400°C. The higher QH measured when T = 200°C is attributed to high dislocation density which is not yet eliminated at 200°C (coldworked steel). For T = 400°C, the high QH-value may be due to G-phase precipitation. This fact requests further investigations. 3.2.3. Recovery effect on the mechanical properties

Kalowing that the duplex steel contains two phases, ferrite and stable austenite, and that these two phases have different hydrogen diffusion coefficients, especially at room temperature (Dr~=~10-9 m= s-~ and EMBRITTLEMENT

~.% -~:~

-

O"

1100

40

50

60

70

80

90

i

i

i

i

i

10

20

30

40

100

1000

[-

K

500

!

0

.

_

.

,

I

50

60

QH600(ppm) Fig. 8. TEM micrograph showing G-phase precipitation on M=3C6 (K) carbides at cs/d grain boundary (after ref. [7]).

Fig. I0. Variation of 2 and QH with the tempering temperature: second C curve.

F. Iacoviello et al./ Materials Science and Engineerb~g A224 (1997) I16-i24

122 100

Oosooc) 90

~

80 '

'~

70

~ ~

(6oo°c) (6so'c) (750"C)

',

~ (750 and 900~C)

(800°C)

~"~2~,-~'~--'<'-" ~---'~-'-'= ~'~

"~,."

~,*.r_r~

Fig. I3. Tensile fracture surface for a 3 h tempering heat treatment at 800°C (charged conditions).

60

50

'~'~---:~2:

t

I

t

v

I

10

20

30

40

50

QH(ppm) Fig. 1i. Variation of 2 with Q~ in the temperature range 6001050°C.

m ~ s-~), when outgassing the hydrogenated steel at room temperature for long periods, for example 3 months, it is then agreed that the ferrite phase will be partially free from hydrogen while the austenitic phase and interfaces (grain boundaries and matrix-precipitated phases) still contain almost the whole hydrogen concentration before outgassing. Outgassing tests conducted on pure ferrite samples charged in the same conditions as duplex samples (200°C, 1.5V per Ag) have shown a hydrogen content of 2 - 3 ppm, confirming the very low hydrogen content of ferrite grains, especially after 3 months at 200C. The hydrogen concentration in the austenitic phase, the trapped hydrogen atoms in the interfaces and the residual hydrogen in the traps inside the ferrite will modify the mechanical behaviour of the duplex steel. Those modifications depend on the microstructure of this steel. Fig. 15 shows that, even after 3 months, the mechanical properties of the duplex steel Z2CND2205 are the same as those measured without outgassing at DH?. ~- 10 - 1 4

room temperature (immediately after hydrogen charging) for all the tempering temperatures, except for the austenitisation temperature (I050°C) where we observe a maximum of recovery (2 = 1). In these conditions, the quantity of hydrogen measured at 600°C under vacuum after outgassing the hydrogenated steel for 3 months is about 18 ppm. This quantity is not sufficient to embrittle the steel heat-treated at I050°C (Fig. 9), where the microstructure is a simple mixture of c~- and 7-phases free from other precipitated phases or carbides. For 800°C, where the volume fraction of G is maximal [10], there is a partial recovery (2 ~ 0.8 instead of 0.7) (Fig. 11). The last result means that the hydrogen atoms are not deeply trapped in the interfaces of g-ferrite or in the grain boundaries. In the temperature range 200-550°C (first C curve), we have also outgassed the samples for 6 months at room temperature and then measured the quantity of hydrogen at 600°C under vacuum. Fig. 16 shows the variation of the residual quantity of hydrogen after outgassing for 6 months at room temperature, 100

/ /

I°

l!

90

80

/ (350-500- 550~C) ~

1

/ 0

(450°c) g

/

0

~

t

l

C

) (4oo°c)

[ [

70

1

/"/ N

60

50 0

Second " C " c u r v e

t

t

r

n

I

10

20

30

40

50

QH(ppm) Fig. I2. Tensile fracture surface for a 3 h tempering heat treatment at 450°C (charged conditions).

Fig. I4. Variation of 2 with QH in the temperatm'e range 200-550°C.

F. Iacoviello et a l . / Materials Science and Eng#zeering A224 (i997) 116-I24 1100 1000 900 800 0" o 700 600 500 400 300 200 600

700

800

900 Rm(~elPa)

1000

110,

Fig. 15. Rm-Tempering temperature curves: curve 1, not under charged conditions (air); curve 2, charged conditions, immediately after charging; and curve 3, charged conditions, after outgassing at room temperature for 3 months. QFi(6 months), and that measured immediately at 600°C, QH(t=0). The difference between QH(t=o) and QH(6 months), i.e., AQH, is aiso plotted. We can observe that at all the temperatures (200-5500C) AQH-values measured exceed 10 ppm. After 3 months outgassing, it is supposed that zXQ~-values are higher than or equaI to those measured after 6 months. Knowing that the ferrite phase is embrittted with only about 2 ppm [22], it is not surprising then that the recovery of the mechanical properties is not possible.

4. Conclusion

In this work, we have shown the sensitivity of the duplex stainless steel Z2CND2205 to hydrogen charged at 200°C for 3 h, for heat-treatment temperatures in the range 200-1050°C. According to the present results, the following conclusions are drawn:

500 o_.,

300

100

,

0

10

,

20

I

,

30

40

50

QH (ppm) Fig. 16. Variation of the hydrogen quantity measured at 600°C under vacuum. QH{,= o~, curve I (immediately after charging); 0~(6 months), curve 2 (after 6 months outgassing at room temperature); AQH, curve 3 ( O H ( t = 0) -- QH(6 months))"

123

(1) The coefficient of hydrogen diffusion, Dt~, has been measured using a thermal technique based on the outgassing under vacuum at several constant temperatures. The results of the DH measurements by the thermal technique are in good agreement with the permeation technique at high temperatures. (2) Considering the duplex stainless steel in asrolled conditions, the outgassed quantity of hydrogen increases with increasing cathodic charging time, also considering charging times higher than the time necessary to saturate the interstitial sites in austenite and in ferrite grains. This is due to the presence of reversible trapping sites, such as dislocations and grain boundaries, and the formation of new trapping sites, such as microvoids. All of these trapping sites are not saturated after the time necessary to saturate the interstitial sites. (3) The mechanical properties, especially Rm, are very sensitive to the tempering temperature, i.e., the microstructure modification. For the higher C curve in the TTT diagram, the relationship between 2 and QH is linear: ~7-phase promotes high embrittlement with low hydrogen quantity, while in a microstructure without precipitated phases (1050°C), a low embrittlement corresponds to a high quantity of hydrogen. For the lower C curve in the TTT diagram, 2 is high (low embrittlemerit), corresponding to QH levels which depend on the presence of high dislocation density (200°C) or probably on the precipitation of G-phase (4000C). (4) The recovery effect on the mechanical properties after outgassing at room temperature for 3 months is not too efficient, for tempering temperatures lower then 700°C. At higher tempering temperatures where a-phase is formed, a partial recovery is observed. A total recovery is obtained for the fully annealed conditions (1050°C). References [1] P. Sentance, in J. Charles and S. Bernhardsson (eds.), Duplex Stainless Steel 9I, Vol. 2, Les Editions de Physique, Beaune, 1991, p. 895. [2] V J. Gadjil, S. Mandziej and B.H. Kolster, in N.R. Moody and A.W. Thompson (eds.), Hydrogen Effects on Materials Behaviour, The Minerals, IVletals and lVlaterials Society, Warrendale, PA, 1990, p. 375. [3] L.J.R. Cohen, J.A. Charles and G.C. Smith, in N.R. IVloody and A.W. Thompson (eds.), Hydrogen Effects o~ Materials Behavzour, VoI. 2, The Minerals, Metals and Materials Society, Warrendale, PA, 1990, p. 263. [4] J.R. Valdez-ValIejo, R.C. Newman and R.P.M. Procter, in N.R. Moody and A.W~ Thompson (eds.); llydrogen Effects on 7~Iaterials Behaviour, The Minerals, Metals and Materials Society, Warrendale, PA, 1990, p. 1003. [5] F. Iacoviello, M. Cavallini, M. Habashi and J. Galland, Stainless Steel 95, Vol. 2, Florence, Italy, 1993, p. 203. [6] M. Guttmann, in N.R. Moody and A.W. Thompson (eds.), Hydrogen Effects on Materials Behaviour, Vol. I, The Minerals, Metals and Materials Society, Warrendale, PA, 1991, p. 79.

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[7] A. Redjajmia, G. Metauer and M. Gantois, in N.R. Moody and A.W. Thompson (eds.), Hydrogen Effects on Materials Behaviour, Vol. 1, The Minerals, Metais and Materials Society, Warrendale, PA, 1991, p. I19. [8] P. Auger, F. Danoix, M. Guttmann and D. Blavette, in N.R. Moody and A.W. Thompson (eds.), Hydrogen Effects on Materials Behaviour, Vol. 1, The Minerals, Metals and Materials Society, Warrendale, PA, 1990, p. 100. [9] F. Danoix, S. Chambreland, J.P. Massoud and P. Auger, in N.R. Moody and A.W. Thompson (eds.), Hydrogen Effects on Materials Behaviour, Vol. 1, The Minerals, Metals and Materials Society, Warrendale, PA, 1990, p. 111. [10] B. Josefsson, J.O. Nilsson and A. Wilson, in N.R. Moody and A.W. Thompson (eds.), Hydrogen Effects on Materials Behaviour, Vol. I, The Minerals. Metals and Matermls Society, Warrendale, PA, 1990, p. 67. [1i] X.G. Wang, D. Dumortier and Y. Riquier. in N.R. Moody and A.W. Thompson (eds.), Hydrogen Effects on MareriaL~ Behaviour, Vol. 1, The Minerals, Metals and Matenals Society, Warrendale, PA, 1990~ p. 127.

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