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Influence of copper as an alloying element on hydrogen environment embrittlement of austenitic stainless steel
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 7 ( 2 0 1 2 ) 1 2 7 6 5 e1 2 7 7 0
Available online at www.sciencedirect.com
journal homepage: www.elsevier.com/locate/he
Technical communication
Influence of copper as an alloying element on hydrogen environment embrittlement of austenitic stainless steel Thorsten Michler a,*, Jo¨rg Naumann b, Erich Sattler c a
Adam Opel AG, 65423 Ruesselsheim, Germany BMW AG, 80788 Munich, Germany c Materialpru¨fanstalt, Stuttgart University, 70511 Stuttgart, Germany b
article info
abstract
Article history:
A Cu alloyed (18Cre10Nie3Cu) and a Cu free (18Cre12.7Ni) austenitic stainless steel were
Received 2 May 2012
tensile tested in gaseous hydrogen atmosphere at 20 C and 50 C. Depending on the test
Received in revised form
temperature, the Cu alloyed steel was extremely embrittled whereas the Cu free steel was
23 May 2012
only slightly embrittled. Austenite stability and inherent deformation mode are two main
Accepted 16 June 2012
criteria for the resistance of austenitic stainless steels against hydrogen environment
Available online 12 July 2012
embrittlement. Based on the well known austenite stability criteria, the austenite stability of both steels should be very similar. Interrupted tensile tests show that martensite
Keywords:
formation upon plastic deformation was much more severe in the Cu alloyed steel
Hydrogen embrittlement
proving that the influence of Cu on austenite stability is overestimated in the empirical
Copper alloyed stainless steel
stability equations. When tested in high pressure H2, replacing Ni by Cu resulted in a fundamental change in fracture mode atmosphere, i.e. Ni cannot be replaced by Cu to reduce the costs of SS without compromising the resistance to hydrogen environment embrittlement. Copyright ª 2012, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved.
1.
Introduction
For automotive hydrogen applications (i.e. fuel cell vehicles), hydrogen environment embrittlement (HEE) is a challenging phenomenon. Due to the well known cost restrictions in automotive industry, safe and cost effective designs are required for hydrogen powered vehicles and materials costs play a major role in alternatively powered cars. HEE and ductile-to-brittle transition (DBT) are well known phenomena in materials science. Common high-pressure compressed hydrogen tanks for automotive applications
operate in a temperature range of 80 to þ85 C and a pressure range of 2e87.5 MPa. To meet both requirements, no HEE and no DBT, CreNi austenitic stainless steels (SS) are common materials for the design of hydrogen wetted components. The best resistance to HEE under severe conditions is reached for Ni contents higher than 12.5 wt% [1,2]. Nickel and Molybdenum are the cost drivers in SS which makes these grades unattractive for automotive applications. Huge cost savings could be reached by replacing Ni by a combination of alternative alloying elements. Within the group of SS austenite stability [3] as well as deformation mode [4] are important
* Corresponding author. E-mail addresses: [email protected] (T. Michler), [email protected] (J. Naumann), [email protected] (E. Sattler). 0360-3199/$ e see front matter Copyright ª 2012, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.ijhydene.2012.06.058
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Table 1 e Chemical compositions, Ms [5] and Md30 [6] temperatures of the Cu-SS (DIN 1.4567, 18Cre10Nie3Cu) and the Ni-SS (18Cre12.7Ni) [11] investigated in this study. Designation
Cu-SS Ni-SS
C
Si
Mn
P
S
Cr
Ni
Cu
N
Mo
wt%
wt%
wt%
wt%
wt%
wt%
wt%
wt%
wt%
wt%
0.011 0.017
0.3 0.5
0.7 1.8
0.024 0.013
0.008 0.007
17.40 18.14
9.8 12.7
3.26 e
0.023 0.017
0.18 0.02
Ms
Md30
C
230 156
C
91 97
Table 2 e Tensile test conditions and results of all tests performed in this study. Steel
Atmosphere
Temp.
Pressure
0.2 Yield Strength
Ultimate Tensile Strength
Elongation at Rupture
Reduction of Area
C
MPa
MPa
MPa
%
%
20 20 20 20 50 50 50 50
10 0.1 10 10 40 0.1 40 10
176 176 183 204 222 209 e 257
502 520 504 516 729 689 712 717
52 60 79 78 48 66 55 77
54 92 84 85 29 82 50 85
Cu-SS Cu-SS Ni-SS Ni-SS Cu-SS Cu-SS Ni-SS Ni-SS
H2 air H2 He H2 air H2 He
criteria for the resistance against HEE. Cu is a candidate element to replace Ni because Cu is one of the elements (among others) stabilizing the austenitic microstructure. Austenite stability is a function of chemical composition. Extensive experimental efforts were made in the 1960’s and 1970’s to establish empirical relationships between chemical composition and austenite stability of metastable SS. Since such equations are easy to use they are still widely used in industry to rank the austenite stability by means of Ms and Md30 temperatures, i.e. martensite start temperature and temperature at which 50% martensite is formed upon a true strain of 30%. The following two equations contain Cu: [5,6] Ms ¼ 502 810C 1230N 13Mn 30Ni 12Cr 54Cu 46Mo
(1)
and Md 30 ¼ 551 462ðC þ NÞ 9; 2Si 8; 1Mn 13; 7Cr 29ðNi þ CuÞ 18; 5Mo 68Nb 1:42 ðGS 8Þ
(2)
with chemical compositions in wt% and GS ¼ ASTM grain size. Both equations show that the impact factor of Cu is equal (Md30) or even higher (Ms) than that of Ni meaning that with respect to austenite stability, 1 wt% Ni should be replaceable roughly by 1 wt% Cu. The second reason for Cu being a candidate element to replace Ni is the initial deformation mode. It was reported in [4] that the susceptibility to HEE is higher in steels with a high degree of inherent slip planarity. Since Cu itself is a wavy glide material [7] and Cu significantly increases the SFE of SS [8], it can be assumed that the influence of Cu alloying on the deformation behaviour of CreNi SS acts in the same way as Ni alloying, i.e. cross slip is facilitated. The metastable austenitic 18Cr-10Ni-3Cu steel (DIN 1.4567) used for this investigation is a commercial Cu alloyed SS. Such grades are characterized by high uniform elongations and less strain hardening [9,10] compared to Cu free steels, e.g. AISI 304. The purpose of this investigation was to assess the susceptibility of the Cu alloyed SS by means of tensile testing in hydrogen atmosphere. The results are compared to a Cu
Fig. 1 e Microstructures in the as-delivered condition. a) Cu-SS, grain size ASTM 7. b) Ni-SS, grain size ASTM 4.
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 7 ( 2 0 1 2 ) 1 2 7 6 5 e1 2 7 7 0
Fig. 2 e Martensite content as a function of true plastic strain. The tests were performed at room temperature and at a constant displacement rate of 0.5 mm/min which corresponds to an engineering strain rate of about 3 * 10L4 1/s. The following parameters were used for the OlseneCohen fit of the Cu-SS: N [ 4.5 (for austenitic stainless steels [18]), a [ 2.8 and b [ 1.2.
free SS of similar Md30 temperature to evaluate the influence of Cu.
2.
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Fig. 4 e Stressestrain diagrams of the tests performed in 10 MPa H2 at 20 C. around 50 to 70 C [12e14], additional testing was performed at 50 C and 40 MPa. Details are given in Table 2. Cylindrical specimens were machined and an engineering strain rate of 5.5 * 105 1/s was used for the tests. The martensite contents of all specimens were measured with a Feritscope MP30E-S by Fischer GmbH, Sindelfingen, Germany. To convert the Feritscope readings into martensite contents the Feritscope readings were multiplied by a factor of 1.7 [15].
Experimental
Steel 1.4567 (18Cr-10Ni-3Cu) and a laboratory SS heat containing 18Cr-12.7Ni [11] were investigated, herein referred to as Cu-SS and Ni-SS, respectively. The chemical compositions of the wrought bar semi finished products are listed in Table 1. Tensile tests in hydrogen (99.9999%) as well as reference tests in air and helium were performed to assess the susceptibility to HEE. The ratio of the reduction of area (RA) in H2 and air (He) was used to quantify the degree of HEE, i.e. RRA ¼ RAH2 =RAairðHeÞ . Baseline testing was performed at room temperature at a gas pressure of 10 MPa. Since it is known that HEE of CreNi austenitic stainless steels reaches a maximum at
3.
Results and discussion
3.1.
Microstructure and austenite stability
The microstructures of both steels investigated here are shown in Fig. 1. Both steels are fully austenitic without any detectable d-ferrite content. The grain sizes are ASTM 7 (CuSS) and ASTM 4 (Ni-SS). To assess the austenite stability as a function of strain, interrupted tensile tests were performed at room temperature
Fig. 3 e Longitudinal cross sections showing the microstructures in the uniform elongation cross sections of the fractured tensile specimens. a) Cu-SS, the dark phase is martensite. b) Ni-SS, the grey lines are mostly slip lines and deformation twins.
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i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 7 ( 2 0 1 2 ) 1 2 7 6 5 e1 2 7 7 0
Fig. 5 e RRA values of the Cu-SS and the Ni-SS tensile tested at 20 C and L50 C.
and martensite contents were measured in defined strain intervals. The data are plotted in Fig. 2 showing the well known sigmoidal shape [16,17] using the OlseneCohen model [18] to fit the data. Massive martensitic transformation was detected in the Cu-SS steel with up to 70% martensite in the necking area of the tensile specimen (which is in fair agreement with the results published in [10]) whereas only 1% martensite was detected in the Ni-SS specimen which is confirmed by the corresponding micrographs (Fig. 3). This result is in total disagreement with the austenite stability predicted by the calculation of the Ms and Md30 temperatures
(Table 1). The 12.5% Ni-SS is known to be very stable with calculated Ms and Md30 temperatures of 156 C and 97 C, respectively. For the Cu-SS the calculated Ms temperature is about 70 C lower compared to the Ni-SS which means that martensitic transformation starts at a lower temperature, i.e. the stability is higher. This cannot be explained by different concentrations of the other alloying elements C, Cr, Mn, Si and Mo because their concentrations are even lower (except Mo) in the Cu-SS. As a consequence, the influence of Cu must be extremely overestimated in [5]. Unfortunately, the data for deriving Eq. (1) is not given in [5] making further interpretations purely speculative. The Md30 temperatures of both Cu-SS and Ni-SS steels are about equal implying a similar martensitic transformation as a function of strain. A Md30 temperature of about 90 C means that 50% of the microstructure is transformed into martensite when 30% strained at a temperature of 90 C. When strained by 30% at room temperature, negligible martensite content was detected in the Ni-SS whereas about 10% martensite was measured in the Cu-SS (Fig. 2) and the difference between both steels increases with increasing strain. Surprisingly, Eq. (2) from Ref. [6] contains a factor for Cu although the steel used for the acquisition of Eq. (2) contained no Cu at all which means that the Cu factor is not justified experimentally. This missing experimental link implies that the effect of Cu was assessed theoretically. The results presented in this study show that the effect of Cu on austenite stability was also overestimated in [6].
3.2.
Hydrogen environment embrittlement tests
The results of the tensile tests in hydrogen and control atmosphere are tabulated in Table 2. One specimen was tested at each condition. The stressestrain curves of the tests
Fig. 6 e Fractographs of the specimens tested in H2. a) Cu-SS tested in 10 MPa H2 at 20 C, b) Cu-SS tested in 40 MPa H2 at L50 C, c) Ni-SS tested in 10 MPa H2 at 20 C, d) Ni-SS tested in 40 MPa H2 at L50 C.
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 7 ( 2 0 1 2 ) 1 2 7 6 5 e1 2 7 7 0
performed in H2 at 20 C are shown in Fig. 4 and the RRA values at 20 C and 50 C indicating the degree of HEE are shown in Fig. 5. The Cu-SS is severely embrittled at room temperature (RRA ¼ 59%) and extremely embrittled at 50 C (RRA ¼ 36%) whereas the Ni-SS is negligibly embrittled at room temperature (RRA ¼ 99%) but severely embrittled at 50 C (RRA ¼ 65%). That is, the degree of embrittlement is higher for the Cu-SS under both test conditions. These results are verified by fractographic analysis (Fig. 6). At both temperatures, 20 and 50 C, respectively, almost the entire fracture cross section of the Cu-SS specimens is characterized by a hydrogen assisted transgranular fracture with elongated voids and facets (Fig. 6a,b) [19,20]. Slip traces are clearly visible on the facets (Fig. 7). For the Ni-SS tested at 20 C, indications of a hydrogen assisted fracture were only observed within a 10e100 mm thick zone at the outer circumference of the fracture surface characterized by a few elongated voids (Fig. 6c). When tested at 50 C, also the Ni-SS showed clear signs of hydrogen assisted fracture (Fig. 6d). There is increasing evidence that the influence of hydrogen on the plastic deformation of stainless steels can be explained by the Hydrogen Enhanced Localized Plasticity (HELP) model. Briefly, hydrogen facilitates planar slip of dislocations by enhancing dislocation motion and suppressing cross slip of dislocations [21] leading to a deformation in bands. Voids nucleate at the intersections of such bands leading to transgranular fracture with different morphologies of elongated voids [19,20]. If slip planarity is promoted by a low stacking fault energy, HELP leads to dislocation pile up at twin or grain boundaries due to enhanced planar slip. Stresses are significantly increased at such obstacles [22] leading to twin or grain boundary separation indicated by facets [19,20] with visible slip traces. With this information and the fractographic results presented in Fig. 6, it can be concluded that under the influence of hydrogen deformation in the Cu-SS is much more localized compared to the Ni-SS. It must be concluded that in SS, Ni cannot be replaced by Cu to reduce the costs of SS without compromising the resistance to HEE.
Fig. 7 e Facet on the fracture surface of the Cu-SS tested in 40 MPa H2 at L50 C. Slip traces are clearly visible.
4.
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Summary
The results can be summarized as follows: According to the well known austenite stability criteria [5,6], austenite stability of the Cu-SS and the Ni-SS investigated in this study should be similar. Interrupted tensile tests showed that martensite formation upon plastic deformation was much more severe in the Cu alloyed steel proving that the influence of Cu on austenite stability is overestimated in the empirical stability equations. The Cu-SS is severely embrittled at 20 C and extremely embrittled at 50 C whereas the Ni-SS is negligibly embrittled at 20 C but severely embrittled at 50 C. Ni cannot be replaced by Cu to reduce the costs of SS without compromising the resistance to HEE.
Acknowledgements This work was partly funded by the German Bundesministerium fu¨r Wirtschaft und Technologie under contract number 0327802A and by the German Stiftung Stahlanwendungsforschung im Stifterverband fu¨r die Deutsche Wissenschaft e.V. under contract number S932.
references
[1] Michler T, Naumann J. Hydrogen environment embrittlement of austenitic stainless steels at low temperatures. Int J Hydrogen Energy 2008;33:2111e22. [2] Michler T, Yukhimchuk A, Naumann J. Hydrogen embrittlement testing at low temperatures and high pressures. Corr Sci 2008;50:3519e26. [3] Martin M, Weber S, Theisen W, Michler T, Naumann J. Effect of alloying elements on hydrogen environment embrittlement of AISI type 304 austenitic stainless steel. Int J Hydrogen Energy 2011;36:15888e98. [4] Thompson AW, Bernstein IM. The role of metallurgical variables in hydrogen-assisted environmental fracture. In: Fontata MG, Staehle RW, editors. Adv in Corr Sci and Technol, 7. Plenum Pub. Corp; 1980. p. 53e175. [5] Pickering FB. Physical metallurgy and the design of steels. London: Appl. Sci. Publ.; 1978. [6] Nohara K, Ono Y, Ohashi N. Strain-induced matensitic transformation in metastabile austenitic stainless steels in multi-stage tensile deformation at various temperatures. In: Proc 1st Int Symp on New Aspects of Martensitic Transformation Kobe, Japan; May 1976. p. 315e20. ¨ ttinger O. History effects in metals [7] Christ HJ, Hoffmann G, O during constant and variable amplitude testing I: wavy dislocation glide behaviour. Mat Sci Eng 1995;A201:1e12. [8] Dulieu D, Nutting J. Influence of solute additions on the stacking fault energy of iron-nickel-chromium austenites. In: Proc. Conf. Metallurgical Developments in High Alloy Steels, Scarborough, UK. The Iron and Steel Inst.; 1964. p. 140e5. [9] Gonzales BM, Castro CSB, Buono VTL, Viela JMC, Andrade MS, Moraes JMD, et al. The influence of copper addition on the formability of AISI 304 stainless steel. Mat Sci Eng 2003;A343:51e6. [10] Choi JY, Jin W. Strain induced martensite formation and its effect on strain hardening behaviour in the cold drawn 304 austenitic stainless steels. Scripta Mater 1997;36:99e104.
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[11] Michler T, Berreth K, Naumann J, Sattler E. Analysis of martensitic transformation in 304 type stainless steels tensile tested in high pressure hydrogen atmosphere by means of XRD and magnetic induction. Int J Hydrogen Energy 2012;37:3567e72. [12] Caskey GR. Hydrogen compatibility handbook for stainless steels. E.I. du Pond de Nemours&Co, Savannah River Lab; 1983. [13] Fukuyama S, Sun D, Zhang L, Wen M, Yokogawa K. Effect of temperature on hydrogen environment embrittlement of type 316 series austenitic stainless steels at low temperatures. J Japan Inst Metals 2003;67:456e9. [14] Han G, He S, Fukuyama S, Yokogawa K. Effect of straininduced martensite on hydrogen environment embrittlement of sensitized austenitic stainless steels at low temperatures. Acta Mater 1998;46:4559e70. [15] Talonen J, Aspegren P, Ha¨nninen H. Comparison of different methods for measuring the strain induced a´-martensite content in austenitic steels. Mat Sci Technol 2004;20: 1506e12. [16] Angel T. Formation of martensite in austenitic stainless steels. J Iron and Steel Inst 1954;5:165e74.
[17] Talonen J, Ha¨nninen H. Formation of shear bands and straininduced martensite during plastic deformation of metastable austenitic steels. Acta Mater 2007;55:6108e18. [18] Olson GB, Cohen M. Kinetics of strain-induced martensitic nucleation. Metall Trans 1975;6A:791e5. [19] San Marchi C, Nibur KA, Balch DK, Somerday BP, Tang X, Schiroky GH, et al. Hydrogen-assisted fracture of austenitic stainless steels. In: Somerday BP, Sofronis P, editors. Hydrogen in metals. Ohio: ASM International, Metals Park; 2009. [20] San Marchi C, Yang NYC, Headley TJ, Michael J. Hydrogenassisted fracture of low nickel content 304 and 316L austenitic stainless steels. Proc 18th European Conf. on Fracture, ECF18, 30Aug e 3Sept 2010, [Dresden, Germany]. [21] Birnbaum HK, Robertson IM, Sofronis P, Teter P. Mechanisms of hydrogen related fracture e a review. In: Magnin T, editor. Proc. corrosion-deformation interactions 96. London: The Institute of Materials; 1997. p. 172e95. [22] Chateau JP, Delafosse D, Magnin T. Numerical simulations of hydrogen-dislocation interactions in fcc stainless steels. Part II: hydrogen effects on crack tip plasticity at a stress corrosion crack. Acta Mater 2002;50:1523e38.
Available online at www.sciencedirect.com
journal homepage: www.elsevier.com/locate/he
Technical communication
Influence of copper as an alloying element on hydrogen environment embrittlement of austenitic stainless steel Thorsten Michler a,*, Jo¨rg Naumann b, Erich Sattler c a
Adam Opel AG, 65423 Ruesselsheim, Germany BMW AG, 80788 Munich, Germany c Materialpru¨fanstalt, Stuttgart University, 70511 Stuttgart, Germany b
article info
abstract
Article history:
A Cu alloyed (18Cre10Nie3Cu) and a Cu free (18Cre12.7Ni) austenitic stainless steel were
Received 2 May 2012
tensile tested in gaseous hydrogen atmosphere at 20 C and 50 C. Depending on the test
Received in revised form
temperature, the Cu alloyed steel was extremely embrittled whereas the Cu free steel was
23 May 2012
only slightly embrittled. Austenite stability and inherent deformation mode are two main
Accepted 16 June 2012
criteria for the resistance of austenitic stainless steels against hydrogen environment
Available online 12 July 2012
embrittlement. Based on the well known austenite stability criteria, the austenite stability of both steels should be very similar. Interrupted tensile tests show that martensite
Keywords:
formation upon plastic deformation was much more severe in the Cu alloyed steel
Hydrogen embrittlement
proving that the influence of Cu on austenite stability is overestimated in the empirical
Copper alloyed stainless steel
stability equations. When tested in high pressure H2, replacing Ni by Cu resulted in a fundamental change in fracture mode atmosphere, i.e. Ni cannot be replaced by Cu to reduce the costs of SS without compromising the resistance to hydrogen environment embrittlement. Copyright ª 2012, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved.
1.
Introduction
For automotive hydrogen applications (i.e. fuel cell vehicles), hydrogen environment embrittlement (HEE) is a challenging phenomenon. Due to the well known cost restrictions in automotive industry, safe and cost effective designs are required for hydrogen powered vehicles and materials costs play a major role in alternatively powered cars. HEE and ductile-to-brittle transition (DBT) are well known phenomena in materials science. Common high-pressure compressed hydrogen tanks for automotive applications
operate in a temperature range of 80 to þ85 C and a pressure range of 2e87.5 MPa. To meet both requirements, no HEE and no DBT, CreNi austenitic stainless steels (SS) are common materials for the design of hydrogen wetted components. The best resistance to HEE under severe conditions is reached for Ni contents higher than 12.5 wt% [1,2]. Nickel and Molybdenum are the cost drivers in SS which makes these grades unattractive for automotive applications. Huge cost savings could be reached by replacing Ni by a combination of alternative alloying elements. Within the group of SS austenite stability [3] as well as deformation mode [4] are important
* Corresponding author. E-mail addresses: [email protected] (T. Michler), [email protected] (J. Naumann), [email protected] (E. Sattler). 0360-3199/$ e see front matter Copyright ª 2012, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.ijhydene.2012.06.058
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i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 7 ( 2 0 1 2 ) 1 2 7 6 5 e1 2 7 7 0
Table 1 e Chemical compositions, Ms [5] and Md30 [6] temperatures of the Cu-SS (DIN 1.4567, 18Cre10Nie3Cu) and the Ni-SS (18Cre12.7Ni) [11] investigated in this study. Designation
Cu-SS Ni-SS
C
Si
Mn
P
S
Cr
Ni
Cu
N
Mo
wt%
wt%
wt%
wt%
wt%
wt%
wt%
wt%
wt%
wt%
0.011 0.017
0.3 0.5
0.7 1.8
0.024 0.013
0.008 0.007
17.40 18.14
9.8 12.7
3.26 e
0.023 0.017
0.18 0.02
Ms
Md30
C
230 156
C
91 97
Table 2 e Tensile test conditions and results of all tests performed in this study. Steel
Atmosphere
Temp.
Pressure
0.2 Yield Strength
Ultimate Tensile Strength
Elongation at Rupture
Reduction of Area
C
MPa
MPa
MPa
%
%
20 20 20 20 50 50 50 50
10 0.1 10 10 40 0.1 40 10
176 176 183 204 222 209 e 257
502 520 504 516 729 689 712 717
52 60 79 78 48 66 55 77
54 92 84 85 29 82 50 85
Cu-SS Cu-SS Ni-SS Ni-SS Cu-SS Cu-SS Ni-SS Ni-SS
H2 air H2 He H2 air H2 He
criteria for the resistance against HEE. Cu is a candidate element to replace Ni because Cu is one of the elements (among others) stabilizing the austenitic microstructure. Austenite stability is a function of chemical composition. Extensive experimental efforts were made in the 1960’s and 1970’s to establish empirical relationships between chemical composition and austenite stability of metastable SS. Since such equations are easy to use they are still widely used in industry to rank the austenite stability by means of Ms and Md30 temperatures, i.e. martensite start temperature and temperature at which 50% martensite is formed upon a true strain of 30%. The following two equations contain Cu: [5,6] Ms ¼ 502 810C 1230N 13Mn 30Ni 12Cr 54Cu 46Mo
(1)
and Md 30 ¼ 551 462ðC þ NÞ 9; 2Si 8; 1Mn 13; 7Cr 29ðNi þ CuÞ 18; 5Mo 68Nb 1:42 ðGS 8Þ
(2)
with chemical compositions in wt% and GS ¼ ASTM grain size. Both equations show that the impact factor of Cu is equal (Md30) or even higher (Ms) than that of Ni meaning that with respect to austenite stability, 1 wt% Ni should be replaceable roughly by 1 wt% Cu. The second reason for Cu being a candidate element to replace Ni is the initial deformation mode. It was reported in [4] that the susceptibility to HEE is higher in steels with a high degree of inherent slip planarity. Since Cu itself is a wavy glide material [7] and Cu significantly increases the SFE of SS [8], it can be assumed that the influence of Cu alloying on the deformation behaviour of CreNi SS acts in the same way as Ni alloying, i.e. cross slip is facilitated. The metastable austenitic 18Cr-10Ni-3Cu steel (DIN 1.4567) used for this investigation is a commercial Cu alloyed SS. Such grades are characterized by high uniform elongations and less strain hardening [9,10] compared to Cu free steels, e.g. AISI 304. The purpose of this investigation was to assess the susceptibility of the Cu alloyed SS by means of tensile testing in hydrogen atmosphere. The results are compared to a Cu
Fig. 1 e Microstructures in the as-delivered condition. a) Cu-SS, grain size ASTM 7. b) Ni-SS, grain size ASTM 4.
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 7 ( 2 0 1 2 ) 1 2 7 6 5 e1 2 7 7 0
Fig. 2 e Martensite content as a function of true plastic strain. The tests were performed at room temperature and at a constant displacement rate of 0.5 mm/min which corresponds to an engineering strain rate of about 3 * 10L4 1/s. The following parameters were used for the OlseneCohen fit of the Cu-SS: N [ 4.5 (for austenitic stainless steels [18]), a [ 2.8 and b [ 1.2.
free SS of similar Md30 temperature to evaluate the influence of Cu.
2.
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Fig. 4 e Stressestrain diagrams of the tests performed in 10 MPa H2 at 20 C. around 50 to 70 C [12e14], additional testing was performed at 50 C and 40 MPa. Details are given in Table 2. Cylindrical specimens were machined and an engineering strain rate of 5.5 * 105 1/s was used for the tests. The martensite contents of all specimens were measured with a Feritscope MP30E-S by Fischer GmbH, Sindelfingen, Germany. To convert the Feritscope readings into martensite contents the Feritscope readings were multiplied by a factor of 1.7 [15].
Experimental
Steel 1.4567 (18Cr-10Ni-3Cu) and a laboratory SS heat containing 18Cr-12.7Ni [11] were investigated, herein referred to as Cu-SS and Ni-SS, respectively. The chemical compositions of the wrought bar semi finished products are listed in Table 1. Tensile tests in hydrogen (99.9999%) as well as reference tests in air and helium were performed to assess the susceptibility to HEE. The ratio of the reduction of area (RA) in H2 and air (He) was used to quantify the degree of HEE, i.e. RRA ¼ RAH2 =RAairðHeÞ . Baseline testing was performed at room temperature at a gas pressure of 10 MPa. Since it is known that HEE of CreNi austenitic stainless steels reaches a maximum at
3.
Results and discussion
3.1.
Microstructure and austenite stability
The microstructures of both steels investigated here are shown in Fig. 1. Both steels are fully austenitic without any detectable d-ferrite content. The grain sizes are ASTM 7 (CuSS) and ASTM 4 (Ni-SS). To assess the austenite stability as a function of strain, interrupted tensile tests were performed at room temperature
Fig. 3 e Longitudinal cross sections showing the microstructures in the uniform elongation cross sections of the fractured tensile specimens. a) Cu-SS, the dark phase is martensite. b) Ni-SS, the grey lines are mostly slip lines and deformation twins.
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Fig. 5 e RRA values of the Cu-SS and the Ni-SS tensile tested at 20 C and L50 C.
and martensite contents were measured in defined strain intervals. The data are plotted in Fig. 2 showing the well known sigmoidal shape [16,17] using the OlseneCohen model [18] to fit the data. Massive martensitic transformation was detected in the Cu-SS steel with up to 70% martensite in the necking area of the tensile specimen (which is in fair agreement with the results published in [10]) whereas only 1% martensite was detected in the Ni-SS specimen which is confirmed by the corresponding micrographs (Fig. 3). This result is in total disagreement with the austenite stability predicted by the calculation of the Ms and Md30 temperatures
(Table 1). The 12.5% Ni-SS is known to be very stable with calculated Ms and Md30 temperatures of 156 C and 97 C, respectively. For the Cu-SS the calculated Ms temperature is about 70 C lower compared to the Ni-SS which means that martensitic transformation starts at a lower temperature, i.e. the stability is higher. This cannot be explained by different concentrations of the other alloying elements C, Cr, Mn, Si and Mo because their concentrations are even lower (except Mo) in the Cu-SS. As a consequence, the influence of Cu must be extremely overestimated in [5]. Unfortunately, the data for deriving Eq. (1) is not given in [5] making further interpretations purely speculative. The Md30 temperatures of both Cu-SS and Ni-SS steels are about equal implying a similar martensitic transformation as a function of strain. A Md30 temperature of about 90 C means that 50% of the microstructure is transformed into martensite when 30% strained at a temperature of 90 C. When strained by 30% at room temperature, negligible martensite content was detected in the Ni-SS whereas about 10% martensite was measured in the Cu-SS (Fig. 2) and the difference between both steels increases with increasing strain. Surprisingly, Eq. (2) from Ref. [6] contains a factor for Cu although the steel used for the acquisition of Eq. (2) contained no Cu at all which means that the Cu factor is not justified experimentally. This missing experimental link implies that the effect of Cu was assessed theoretically. The results presented in this study show that the effect of Cu on austenite stability was also overestimated in [6].
3.2.
Hydrogen environment embrittlement tests
The results of the tensile tests in hydrogen and control atmosphere are tabulated in Table 2. One specimen was tested at each condition. The stressestrain curves of the tests
Fig. 6 e Fractographs of the specimens tested in H2. a) Cu-SS tested in 10 MPa H2 at 20 C, b) Cu-SS tested in 40 MPa H2 at L50 C, c) Ni-SS tested in 10 MPa H2 at 20 C, d) Ni-SS tested in 40 MPa H2 at L50 C.
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performed in H2 at 20 C are shown in Fig. 4 and the RRA values at 20 C and 50 C indicating the degree of HEE are shown in Fig. 5. The Cu-SS is severely embrittled at room temperature (RRA ¼ 59%) and extremely embrittled at 50 C (RRA ¼ 36%) whereas the Ni-SS is negligibly embrittled at room temperature (RRA ¼ 99%) but severely embrittled at 50 C (RRA ¼ 65%). That is, the degree of embrittlement is higher for the Cu-SS under both test conditions. These results are verified by fractographic analysis (Fig. 6). At both temperatures, 20 and 50 C, respectively, almost the entire fracture cross section of the Cu-SS specimens is characterized by a hydrogen assisted transgranular fracture with elongated voids and facets (Fig. 6a,b) [19,20]. Slip traces are clearly visible on the facets (Fig. 7). For the Ni-SS tested at 20 C, indications of a hydrogen assisted fracture were only observed within a 10e100 mm thick zone at the outer circumference of the fracture surface characterized by a few elongated voids (Fig. 6c). When tested at 50 C, also the Ni-SS showed clear signs of hydrogen assisted fracture (Fig. 6d). There is increasing evidence that the influence of hydrogen on the plastic deformation of stainless steels can be explained by the Hydrogen Enhanced Localized Plasticity (HELP) model. Briefly, hydrogen facilitates planar slip of dislocations by enhancing dislocation motion and suppressing cross slip of dislocations [21] leading to a deformation in bands. Voids nucleate at the intersections of such bands leading to transgranular fracture with different morphologies of elongated voids [19,20]. If slip planarity is promoted by a low stacking fault energy, HELP leads to dislocation pile up at twin or grain boundaries due to enhanced planar slip. Stresses are significantly increased at such obstacles [22] leading to twin or grain boundary separation indicated by facets [19,20] with visible slip traces. With this information and the fractographic results presented in Fig. 6, it can be concluded that under the influence of hydrogen deformation in the Cu-SS is much more localized compared to the Ni-SS. It must be concluded that in SS, Ni cannot be replaced by Cu to reduce the costs of SS without compromising the resistance to HEE.
Fig. 7 e Facet on the fracture surface of the Cu-SS tested in 40 MPa H2 at L50 C. Slip traces are clearly visible.
4.
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Summary
The results can be summarized as follows: According to the well known austenite stability criteria [5,6], austenite stability of the Cu-SS and the Ni-SS investigated in this study should be similar. Interrupted tensile tests showed that martensite formation upon plastic deformation was much more severe in the Cu alloyed steel proving that the influence of Cu on austenite stability is overestimated in the empirical stability equations. The Cu-SS is severely embrittled at 20 C and extremely embrittled at 50 C whereas the Ni-SS is negligibly embrittled at 20 C but severely embrittled at 50 C. Ni cannot be replaced by Cu to reduce the costs of SS without compromising the resistance to HEE.
Acknowledgements This work was partly funded by the German Bundesministerium fu¨r Wirtschaft und Technologie under contract number 0327802A and by the German Stiftung Stahlanwendungsforschung im Stifterverband fu¨r die Deutsche Wissenschaft e.V. under contract number S932.
references
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