The stress corrosion cracking behavior of austenitic stainless steels in boiling magnesium chloride solutions

The stress corrosion cracking behavior of austenitic stainless steels in boiling magnesium chloride solutions

Corrosion Science 49 (2007) 3040–3051 www.elsevier.com/locate/corsci The stress corrosion cracking behavior of austenitic stainless steels in boiling...

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Corrosion Science 49 (2007) 3040–3051 www.elsevier.com/locate/corsci

The stress corrosion cracking behavior of austenitic stainless steels in boiling magnesium chloride solutions Osama M. Alyousif a b

a,*

, Rokuro Nishimura

b

Department of Mechanical Engineering, Kuwait University, P.O. Box 5969, Safat 13060, Kuwait Department of Applied Materials Science, 1-1, Gakuen-cho, Sakai, Osaka Prefecture University, Osaka 599-8531, Japan Received 30 June 2006; accepted 25 December 2006 Available online 4 March 2007

Abstract The change in the mechanism of stress corrosion cracking with test temperature for Type 304, 310 and 316 austenitic stainless steels was investigated in boiling saturated magnesium chloride solutions using a constant load method. Three parameters (time to failure; tf, steady-state elongation rate; lss and transition time at which a linear increase in elongation starts to deviate; tss) obtained from the corrosion elongation curve showed clearly three regions; stress-dominated, stress corrosion crackingdominated and corrosion-dominated regions. In the stress corrosion cracking-dominated region the fracture mode of type 304 and 316 steels was transgranular at higher temperatures of 416 and 428 K, respectively, but was intergranular at a lower temperature of 408 K. Type 310 steel showed no intergranular fracture but only transgranular fracture. The relationship between log lss and log tf for three steels became good straight lines irrespective of applied stress. The slope depended upon fracture mode; 2 for transgranular mode and 1 for intergranular mode. On the basis of the results obtained, it was estimated that intergranular cracking was resulted from hydrogen embrittlement due to strain-induced formation of martensite along the grain boundaries, while transgranular cracking took place by propagating cracks nucleated at slip steps by dissolution.  2007 Elsevier Ltd. All rights reserved. Keywords: A. Austenitic stainless steels; A. Magnesium chloride solution; C. Steady-state elongation rates; C. Stress corrosion cracking; C. Hydrogen embrittlement

*

Corresponding author. Tel.: +965 498 5793. E-mail address: [email protected] (O.M. Alyousif).

0010-938X/$ - see front matter  2007 Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2006.12.023

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1. Introduction The stress corrosion cracking behavior of austenitic stainless steels in chloride and other corrosive solutions has been extensively investigated using various methods [1–7]. Stress corrosion cracking in the broad sense for austenitic stainless steels is caused by either stress corrosion cracking in the narrow sense (denoted as SCC hereafter) such as active path dissolution [1,2] and film rupture [3,4], or hydrogen embrittlement (denoted as HE) [5–7]. Austenitic stainless steels such as type 304 and 316 steels, but not type 310 steel, are known to undergo phase transformation from c ! a 0 martensite due to applied stress or hydrogen charging [8,9] and a 0 -martensite is considered to be directly related to brittle fracture [10]. However, the role of martensite in SCC and HE is still not fully identified. In a previous paper, it was found that austenitic stainless steels suffered from cracking failure by two different mechanisms of SCC and HE depending upon test temperature [11]. In this study, the failure characteristics of type 304, 316 and 310 austenitic stainless steels were examined under constant loads in boiling magnesium chloride solutions of different temperatures to evaluate the role of martensite in determining the cracking mechanism of austenitic stainless steels. 2. Experimental The specimens used were the commercial type 304, 316 and 310 austenitic stainless steels whose chemical compositions (wt%) are shown in Table 1. As shown in Fig. 1, the geometry for stress corrosion cracking experiments is as follows: the gauge length is 25.6 mm, the width 5 mm and the thickness 1 mm. The specimens were solution-annealed at 1373 K for 3.6 ks under an argon atmosphere and then water-quenched. Prior to experiments, the solution-annealed specimens were polished to 1000 grit emery paper, degreased with acetone in an ultrasonic cleaner and washed with distilled water. After the pretreatment, the specimens were immediately set into a stress corrosion cracking cell. Stress corrosion cracking tests were conducted in boiling saturated magnesium chloride solutions, Table 1 Chemical compositions (wt%) and mechanical properties of the austenitic stainless steels used

SS304 SS310 SS316

C

Si

Mn

P

S

Ni

Cr

Mo

rYield (MPa)

rTensile (MPa)

0.06 0.05 0.06

0.35 0.84 0.70

0.96 1.27 0.96

0.027 0.016 0.031

0.004 0.001 0.004

8.13 19.30 10.15

18.20 24.76 16.98

– – 2.22

276 275 323

691 575 636

Fig. 1. Geometry of specimens (dimensions in mm).

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whose boiling temperatures were changed by the change in the concentration of magnesium chloride. The test temperatures used were 428 K (type 316), 416 K (type 304 and 310) and 408 K (type 304 and 316), where cracking failure was observed under constant applied stress condition [11]. The applied stresses were in the range from 0 to 500 MPa. All experiments were carried out under an open circuit condition. A lever-type constant load apparatus (lever ratio 1:10) to which three specimens can be separately and simultaneously attached was used with a cooling system on the top of a testing cell to avoid evaporation of the solution during the experiments. The specimens were insulated from rod and grip by surface-oxidized zirconium tube. Elongation of the specimens under a constant load was measured by an inductive linear transducer with an accuracy of ±0.01 mm. 3. Results 3.1. Corrosion elongation curve Fig. 2 shows the corrosion elongation curves for type 304 steel at 408 K and type 310 steel at 416 K under a constant applied stress condition (r = 350 MPa) in boiling saturated magnesium chloride solutions, where the corrosion elongation curve of type 304 steel at 408 K was similar to that of type 316 steel at 408 K. From these curves the three parameters were obtained for each specimen: the time to failure (tf), the steady-state elongation rate (lss) for the straight part of the corrosion elongation curve and the ratio of transition time to time to failure (tss/tf), where the transition time (tss) is the time when the elongation curve begins to deviate from the linear increase. For type 304 steel at 408 K, tf was considerably shorter than for type 310 steel at 416 K and tss was close to tf. Thus the tss/tf 10

Boiling Saturated MgCl2 (σ= 350MPa)

SS 310 at T=416K

Displacement / mm

8

6

lss SS 304 at T=408K

4

lss 2

tf

t ss

0 0

2000

4000

tf

t ss 6000

8000

10000

12000

Time / s Fig. 2. Corrosion elongation curves for type 304 steel at 408 K and type 310 steel at 416 K under a constant applied stress condition (r = 350 MPa) in boiling saturated MgCl2 solutions.

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value was close to unity. For type 310 steel at 416 K, tss was approximately half the time to failure. 3.2. Change in three parameters (tf, lss and tss/tf) with solution temperature Fig. 3 depicts the relationships between applied stress (r) and logarithm of time to failure (tf) for the three stainless steels at three test temperatures used. All relationships can be clearly divided into three regions shown by Arabic numerals 1–3 in Fig. 3; region 1 is the stress-dominated region, region 2 the stress corrosion cracking-dominated region and region 3 the corrosion-dominated region. In region 2, tf values for type 304 steel at 416 K and for type 316 steel at 428 K are shorter than those at 408 K and tf for type 304 steel is shorter than that for type 316 steel. In addition, the applied stress range in region 2 for type 316 steel at 408 K is narrower than that for type 304 steel at 408 K; the maximum applied stress in region 2 is the same for both steels, but the minimum applied stress for type 316 steel is higher than that for type 304 steel. At higher temperatures, the applied stress range in region 2 is almost the same for three stainless steels, and the time to failure is in the order of type 304 < type 316 < type 310. The relationships between applied stress (r) and logarithm of steady-state elongation rate (lss) for all stainless steels are shown in Fig. 4. These relationships can be divided into three regions with the same applied stress ranges as those in Fig. 3. The elongation rate at 408 K is almost one order of magnitude lower than those at higher temperatures. The relationships between applied stress (r) and transition time to time to failure ratio (tss/tf) for all stainless steels are shown in Fig. 5. The values of tss/tf are nearly constant in region 2. For type 316 and 304 steels, the value of tss/tf is divided into two groups. In the first group at higher temperatures tss/tf ffi 0.57 ± 0.02 as well as for type 310 steel and in the second group at 408 K tss/tf = 0.85–0.90. 3.3. A parameter for predicting time to failure Fig. 6 shows the relationships between logarithm of steady-state elongation rate (lss) in region 2 in Fig. 4 and logarithm of time to failure (tf) in region 2 in Fig. 3 for all stainless steels. It is noteworthy that the log lss–log tf relation becomes straight lines depending upon material and test temperature. Type 304 steel at 416 K and type 316 steel at 428 K show a common single line with a slope of m = 2. The slope for type 310 is also 2. By contrast, at 408 K type 304 and 316 steels exhibit two straight lines with the same slope of m = 1 as each other. Thus, the empirical equations for the stainless steels are as follows: For type 304 and 316 steels at higher temperatures and type 310 steel in region 2 log lss ¼ 2 log tf þ C 1

ð1Þ

For type 304 and 316 steels in region 2 at 408 K log lss ¼  log tf þ C 2

ð2Þ

where C1 and C2 are constants that depend on the material. It can be seen in Fig. 6 that type 304 steel at 416 K and type 316 steel at 428 K show the same C1 value, although type 310 steel shows higher C1 value, and that the C2 value of type 304 steel at 408 K is higher than that of type 316 steel at 408 K.

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500

310-416K

1

Stress(σ)/MPa

400

300

2

200

100 3

0 500

1

304-416K 304-408K

1

Stress(σ) / MPa

400

300

2

2

3

3

200

100

0 500

1

1

316-428K 316-408K

Stress(σ) / MPa

400 2

300 2

200 3

100 3

0

0

104 105 103 Time to Failure (tf) / s

106

Fig. 3. The relationship between applied stress (r) and logarithm of time to failure (tf) for type 310, 304 and 316 steels in boiling saturated MgCl2 solutions.

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500

3045

1

Stress (σ) / MPa

400

300 2

200

100 3

310- 416K

0 1

Stress (σ) / MPa

500

1

400

300

2

2

200

100

3

304-416K 304-408K

3

0

Stress (σ) / MPa

500

1

1

400 2

300

2

200 3

100 316-428K 316-408K

3

0

10-10

10-9

10-8

10-7

10-6

10-5

Elongation Rate (lss ) / m.s-1 Fig. 4. The relationship between applied stress (r) and logarithm of steady-state elongation rate (lss) for type 310, 304 and 316 steels in boiling saturated MgCl2 solutions.

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Fig. 5. The relationship between applied stress (r) and tss/tf for type 310, 304 and 316 steels in boiling saturated MgCl2 solutions.

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Elongation Rate (lss) / m.s-1

10-5

10-6

316-428K 316-408K 304-416K 304-408K 310-416K

m = -2

m = -2

10-7

10-8

m = -1

Boiling Saturated MgCl2

10-9 102

103

104

105

106

Time to Failure (tf) / s Fig. 6. The relationship between logarithms of steady-state elongation rate (lss) and time to failure (tf) for type 310, 304 and 316 steels at 408, 416 and 428 K in boiling saturated MgCl2 solutions, where m is the slope of the straight line.

In addition to these characteristics, it can be said that steady-state elongation rate (lss) is a useful parameter for predicting the time to failure (tf) independent of applied stress even in boiling saturated magnesium chloride solutions as well as in hydrochloric and sulfuric acid solutions [13,14], because lss can be obtained at a time within 10–20% of tf from the corrosion elongation curve. 3.4. Fracture appearance The fracture appearances of the austenitic stainless steels were observed by scanning electron microscopy. Fig. 7 shows the fracture appearances of type 304 steel (a) and 310 steel (b) in region 2 at 416 K. The fracture mode for type 304 and 310 steels is

Fig. 7. Fracture appearances for type 304 steel at 416 K and r = 300 MPa; ·550 (a) and for type 310 steel at 416 K and r = 300 MPa; ·650 (b) in boiling saturated MgCl2 solution.

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Fig. 8. Fracture appearances for type 316 steel at 408 K and r = 300 MPa; ·550 (a) and for type 316 steel at 408 K and r = 400 MPa; ·1000 (b) in a boiling saturated MgCl2 solution.

predominantly transgranular. The transgranular fracture was observed over the whole applied stresses in region 2. Type 316 steel in region 2 at 428 K exhibited the same transgranular fracture appearance as those in Fig. 7. On the other hand, as shown in Fig. 8, the fracture appearance in region 2 for type 316 steel at 408 K is a mixture of intergranular and transgranular modes (a) at a low applied stress (300 MPa) and predominantly intergranular mode (b) at a high applied stress (400 MPa). The ratio of intergranular appearance to transgranular one increased with increasing applied stress in region 2 and finally the fracture mode became completely intergranular. The fracture appearances for type 304 steel at 408 K were mostly intergranular and almost the same as those for type 316 steel at 408 K showing an increase in the ratio of intergranular appearance to transgranular one with increasing applied stress. For type 310 steel, the fracture mode was transgranular over the whole applied stresses in region 2. 4. Discussion 4.1. Characteristics of stress corrosion cracking In the applied stress region where the stress corrosion cracking-dominated fracture occurred, the fracture mode for type 304 and 316 steels at 416 and 428 K was transgranular, while that at 408 K was mostly intergranular, in particular at higher applied stresses. In other words, for type 304 and 316 steels in boiling magnesium chloride solutions, the fracture mode was transgranular at higher temperatures and intergranular at the lower temperature. The minimum applied stress to induce intergranular cracking for type 316 steel was significantly higher than that for type 304. For type 310 steel only transgranular fracture was found as far as stress corrosion cracking occurred in any temperatures in boiling magnesium chloride solutions [11]. In general, transgranular cracking was characterized by a higher steady-state elongation rate in both stress corrosion cracking-dominated and stress-dominated regions and shorter time to failure. By contrast the failure in intergranular cracking was slow and most of the time to failure was in the steady-state elongation region indicating the small

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mechanical elongation until failure. Consequently, intergranular cracking for these austenitic stainless steels in the boiling magnesium chloride solution revealed more brittle nature in comparison with transgranular cracking. 4.2. A qualitative explanation of fracture mechanism Type 304 steel is a metastable austenite and is highly susceptible to martensite formation [16]. Type 316 steel is also a metastable austenite and is moderately susceptible to martensite formation [17]. On the other hand, type 310 steel is known to be less susceptible to martensite formation [18]. For metastable austenitic stainless steels like type 304 and 316 steels, the strain-induced formation of martensite tends to take place along grain boundaries and particularly is facilitated by hydrogen entry [19,20]. The presence of martensite is known to induce hydrogen embrittlement (HE) because of very high hydrogen diffusivity coefficient and very small hydrogen content compared to those of the austenite [17]. Thus, martensite formed at grain boundaries is apt to be responsible for HE of intergranular mode, and type 304 and 316 steels have higher susceptibility to martensite-induced HE of intergranular mode, but type 310 has little susceptibility to martensite-induced HE of intergranular mode. On the other hand, transgranular cracking is caused by propagation of cracks nucleated at slip steps and is not related to martensite formed at grain boundaries. Thus stress corrosion cracking in the narrow sense of the word (SCC) includes such transgranular cracking. Type 304 and 316 steels suffered intergranular cracking at a lower temperature (408 K) and transgranular cracking at higher temperatures (428 K for type 316 and 416 K for type 304). In determining the cracking mode a competition between dissolution of material at slip steps inducing SCC and hydrogen entry inducing HE is the decisive factor as well as the formation of martensite. The hydrogen entry is determined by the difference between hydrogen absorption and hydrogen escape. Transgranular SCC for type 304 and 316 steels was caused by propagation of cracks nucleated at slip steps, not at martensite at grain boundaries [21]. Therefore, the cracking mode would be determined by the competition between the dissolution rate at slip steps and the hydrogen entry rate at grain boundaries with martensite. As test temperature increases, the dissolution rate increases. On the other hand, both hydrogen absorption rate and hydrogen escape rate increase with increasing test temperature, whereas the amount of hydrogen entry decreases, because the hydrogen escape rate becomes superior to the hydrogen entry rate with increasing temperature. This means that the amount of hydrogen entry decreases with increasing test temperature [22]. In addition, it is considered that the amount of martensite decreases with increasing test temperature, which suggests that the hydrogen entry rate decreases with increasing test temperature [23]. Thus, the dissolution rate becomes higher than the hydrogen entry rate at higher temperatures. This will result in transgranular SCC at higher temperatures. On the other hand, it is known that the amount of martensite increases with strain or applied stress [17]. At 408 K for type 304 and 316 steels, the fracture mode changed from a mixed transgranular and intergranular mode to a complete intergranular mode with increasing applied stress. If cracking failure at 408 K is caused by martensite-induced HE, the increase in intergranular mode with increasing applied stress is in agreement with an increase in the amount of martensite with increasing applied stress. Similarly, the minimum applied stress to induce intergranular cracking for type 316 steel was significantly

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higher than that for type 304 steel in agreement with the fact that type 316 steel is less susceptible to stress-induced martensite formation than type 304 steel. Furthermore, type 304 steel was more susceptible to intergranular cracking in comparison with type 316 steel in agreement with the fact the 304 steel has higher susceptibility to strain-induced martensite formation in comparison with type 316 steel. Consequently, intergranular cracking observed at the lower temperature, that is, 408 K for type 304 and 316 steels is due to grain boundary martensite-induced hydrogen embrittlement (HE) and transgranular cracking for type 304 and 316 steels at higher temperatures at 416 and 428 K and for 310 steel is ascribed to stress corrosion cracking in the narrow sense of the word (SCC). 5. Conclusions The stress corrosion cracking behavior of three austenite stainless steels with different susceptibilities to stress-induced martensite formation was examined in boiling saturated magnesium chloride solutions with different boiling temperatures under constant load condition. The following conclusions can be drawn: (1) The applied stress dependences of the three parameters time to failure; tf, steadystate elongation rate;lss and transition time at which a linear increase in elongation starts to deviate; tss obtained from the corrosion elongation curve clearly showed three regions; the stress-dominated region, the stress corrosion cracking-dominated region and the corrosion-dominated region. (2) The fracture mode in the stress corrosion cracking-dominated region for type 304 and 316 steels was transgranular at higher temperatures of 416 and 428 K and was intergranular at a lower temperature of 408 K. Type 310 steel showed no intergranular fracture but only transgranular fracture. (3) The time to failure (tf) mostly by transgranular cracking was in the order of type 304 < type 316 < type 310 and that by intergranular cracking for type 304 steel was shorter than type 316 steel. (4) The relationships between log lss and log tf for three steels showed good straight lines irrespective of applied stresses and test temperatures. However, the slope of the linear relationship depended upon fracture mode; 2 for transgranular mode and 1 for intergranular mode. (5) Intergranular cracking was attributable to hydrogen embrittlement (HE) caused by strain-induced formation of martensite along grain boundaries with hydrogen entry, while transgranular cracking was due to active path corrosion nucleated at slip steps by dissolution. References [1] [2] [3] [4] [5] [6] [7]

T.P. Hoar, J.M. West, Proceeding of the Royal Society A268 (1962) 304–315. T.P. Hoar, J.C. Scully, Journal of Electrochemistry 111 (1964) 348–352. R. Nishimura, Y. Maeda, Corrosion Science 46 (2004) 343–355. T. Nakayama, M. Takano, Corrosion 42 (1986) 10–14. M.B. Whiteman, A.R. Troiano, Corrosion 21 (1965) 53–56. M.L. Holzworth, Corrosion 25 (1969) 107–115. P.R. Rhodes, Corrosion 25 (1969) 462–467.

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[8] H. Hanninen, T. Hakkarainen, Metallurgical Transactions A 10A (1979) 1196–1199. [9] N. Narita, C.J. Altstetter, H.K. Birnbaum, Metallurgical Transactions A 13A (1982) 1355–1365. [10] J.C. Scully, The role of hydrogen in stress corrosion cracking, in: A.W. Thompson, I.M. Bernstein (Eds.), Effect of Hydrogen on Behavior of Materials, AIME, 1975, p. 129. [11] O.M. Alyousif, R. Nishimura, Corrosion Science 48 (2006) 4283–4293. [13] R. Nishimura, K. Kudo, Corrosion 45 (1989) 308–316. [14] R. Nishimura, Corrosion Science 34 (1993) 1859–1868. [16] C.L. Briant, Metallurgical Transactions A 10A (1979) 181–189. [17] X. Sun, J. Xu, Y. Li, Acta Metallurgica 37 (1989) 2171–2176. [18] S.S. Birley, D. Tromans, Corrosion 27 (1971) 63–71. [19] T-P. Perng, C.J. Altstetter, Acta Metallurgica 34 (1986) 1771–1781. [20] H. Hanninen, T. Hakarainen, Corrosion 36 (1980) 47–51. [21] R. Nishimura, Corrosion 46 (1990) 311–318. [22] J.K. Tien, Diffusion and the dislocation sweeping mechanism, in: A.W. Thompson, I.M. Bernstein (Eds.), Effect of Hydrogen on Behavior of Materials, AIME, 1975, p. 309. [23] R.E. Reed-Hill, Physical Metallurgy Principles, second ed., Brooks/Cole Engineering Division, 1973, p. 656.

Further reading [12] R. Nishimura, A. Sulaiman, Y. Maeda, Corrosion Science 45 (2003) 465–484. [15] R. Nishimura, Corrosion Science 34 (1993) 1463–1473.