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Hydrogen environment embrittlement of an ODS RAF steel – Role of irreversible hydrogen trap sites
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 5 ( 2 0 1 0 ) 9 7 4 6 e9 7 5 4
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Hydrogen environment embrittlement of an ODS RAF steel e Role of irreversible hydrogen trap sites Thorsten Michler a,*, Michael P. Balogh b a b
Adam Opel GmbH, Ru¨sselsheim, Germany General Motors Research and Development, Warren, MI, USA
article info
abstract
Article history:
The microstructure and the effects of 10 MPa hydrogen atmosphere on the tensile prop-
Received 3 May 2010
erties of a oxide dispersion strengthened (ODS) reduced activation ferritic (RAF) steel were
Received in revised form
investigated. The microstructure consists of a fine grained ferritic matrix with Me3O4
24 June 2010
(Me ¼ Cr, Fe or Mn), VN and Cr23C6 grain boundary precipitates as well as dispersed yttrium
Accepted 24 June 2010
oxide nano precipitates in the ferritic matrix. The yield and ultimate tensile strength were
Available online 1 August 2010
unaffected by the H2 atmosphere whereas elongation at fracture and reduction in area were markedly reduced. In H2 atmosphere, the fracture morphology was found to be
Keywords:
a mixture of intergranular H-assisted fracture and a smaller amount of transgranular
Hydrogen
hydrogen enhanced localized plasticity (HELP) fracture. The sensitivity of the ODS RAF steel
Hydrogen traps
to hydrogen embrittlement is attributed to the large number grain boundary precipitates
Hydrogen embrittlement
which enhance the tendency for intergranular fracture.
ODS steel
ª 2010 Professor T. Nejat Veziroglu. Published by Elsevier Ltd. All rights reserved.
Slow displacement rate tensile test
1.
Introduction
Hydrogen environment embrittlement (HEE) is a well known phenomenon 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 20e875 bar. Common materials for the design of hydrogen wetted components are CreNi austenitic stainless steels (SS). For stationary hydrogen tanks where cost and weight are of minor importance, 18Cr-10Ni-2Mo steels like DIN 1.4404 (AISI 316L) or DIN 1.4571 (AISI 316 Ti) are widely and successfully used. The best resistance to HEE under severe conditions is reached for Ni contents higher than 12.5 wt% in solution treated condition [1,2]. Nickel and molybdenum are the cost drivers in SS which makes these grades unattractive for automotive mass production. Furthermore, the strength of SS is quite low (yield strength from about 200 to 250 MPa)
resulting in thick tank walls and ultimately making such tanks not only expensive but also quite heavy. Significant cost and weight savings would be possible if a higher strength steel could be utilized. Unfortunately, for conventional steels, hydrogen enhances the susceptibility of high strength steels to brittle fracture, i.e. the higher the strength the higher the susceptibility [3] in addition to the fact that high strength steels are intrinsically more prone to brittle cracking than lower strength steels. Today, the most accepted explanation for HEE is the hydrogen enhanced localized plasticity (HELP) model. According to HELP, dissolved H enhances dislocation mobility and slip planarity [4e7], which can lead to heterogeneously localized plastic strain and stress concentrations presumably sufficient to enable subcritical crack growth. Due to enhanced-heterogeneous planar slip with H present, the macroscopic fracture appearance is “brittle” and involves
* Corresponding author. E-mail address: [email protected] (T. Michler). 0360-3199/$ e see front matter ª 2010 Professor T. Nejat Veziroglu. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.ijhydene.2010.06.071
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Table 1 e Some reversible and irreversible traps and their binding energies relative to atomic hydrogen. Trap
Binding energy Reference [kJ/mol]
C interstitial Interstitial site in iron based alloy Dislocation in Fe Screw dislocation core Grain boundary Fe oxide interface Y2O3 interface Al2O3 interface Cr carbide interface Fe3C interface
3 4e8 24e26 20e30 18e53 47 70 79 67 84
[12] [12] [11,12] [11] [11] [11] [11] [11] [13] [11]
either slip band interface cracking or microcrack formation at slip band, twin and/or grain boundary intersections [8e10]. Another characteristic of hydrogen in steels is that hydrogen is not distributed homogeneously in the lattice. In general, hydrogen is trapped at higher binding energy
Table 2 e Chemical composition.
C Si Mn Cr W V Ta Y 2O 3
Nom. composition, wt %
Meas. composition, wt %
0.1 0.1 0.4 9 1.1 0.2 0.15 0.3e0.5
Not measured 0.05 0.38 8.7 0.66 0.19 0.04 0.25 (0.22 Y)
sites in comparison to normal interstitial sites. Trapping sites fall into two classes depending on the binding energy. Traps with binding energies less than w50 kJ/mol are called reversible traps, whereas those with higher energies are termed irreversible traps. Reversible traps are characterized by the ability to capture and release hydrogen atoms. Due to their high binding energies, irreversible traps only
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capture hydrogen until they become saturated and release of hydrogen atoms is very difficult if not impossible. As shown in Table 1, low binding energies are reported for common structural features like grain boundaries, dislocations or substitutional elements while high binding energies are found for vacancies and precipitates like carbides and oxides [11e13]. Recent studies on carbidic precipitates have shown that the trapping capacity also depends on the lattice mismatch, i.e. whether the carbides are coherent or incoherent [14,15]. The role of trapping upon hydrogen embrittlement is not clear yet. Trapping might have a positive effect since the amount of diffusible hydrogen is strongly reduced. On the other hand, trapping might be detrimental when hydrogen weakens the structural integrity, e.g. at grain boundaries. The concept of incorporating irreversible H traps into the lattice to reduce HEE was proposed quite some time ago (see review [16] and references therein) but systematic investigations are missing. To reduce the amount of diffusible hydrogen in the lattice, traps need to be homogenously distributed, have a high binding energy, available in sufficient quantity and separated by a minimal distance [16]. IF HELP is assumed to be the dominant mechanism, the effect of H on dislocation mobility might be reduced when H is trapped in irreversible traps. Since the mobility of H in HEE experiments at room temperature is only a few mm, fine dispersed nano precipitates could be promising. Reduced activation ferritic/martensitic (RAFM) steels are materials for structural applications in future fusion reactors. The main requirements are a high resistance to both transmutation He-induced embrittlement as well as to neutron irradiation embrittlement which decreases the fracture toughness of many materials and causes an increase in ductile-to-brittle transition temperature. One candidate material is EUROFER 97 with a base composition of 0.1 wt% C, 9 wt% Cr, 0.4 wt% Mn and 0.1 wt% Si. Other (micro) alloying elements are 1.1 wt% W, 0.2 wt% V and 0.15 wt% Ta added to reduce neutron radiation induced swelling. About 9 wt% Cr is added to reduce irradiation hardening. Furthermore, the superior performance of RAFM steels is attributed to the high helium trapping capacity of the structure, which contains martensite laths, precipitates and carbides as traps with high binding energies [17,18]. To increase the operating temperature and thus the efficiency of the reactors, mechanical
Fig. 1 e Macroscopic structure of the ODS EUROFER steel. a) Light micrograph V2A etchant, b). SEM image, the vertical structures are artifacts from the FIB polishing process.
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Fig. 2 e SEM EDS elemental maps collected from the FIB polished specimen, see Fig. 1b. a) Fe map, b) Cr map, c) Mn map, d) V map.
milling is used to embed finely dispersed oxide particles (e.g. Y2O3) into the base alloy (e.g. EUROFER 97) which increases the creep strength [19,20]. Over the past years, oxide dispersion strengthened (ODS) alloys have attracted much attention and
their properties have been characterized in detail [17e19,21e24]. Due to Y2O3 alloying, the structure is entirely ferritic [19] and thus, such steels are denoted reduced activation ferritic (RAF) steels.
Fig. 3 e EFTEM images of a large oxide precipitate. a) bright field image, b) CrL map, c) MnL map, d) VL map, e) OK map, f) NK map.
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Fig. 4 e High resolution TEM image of a large oxide precipitate.
Hydrogen trapping and transport in ODS-EUROFER was investigated in [25] showing that Y2O3 precipitates act as irreversible hydrogen traps. Severe influence of hydrogen upon tensile ductility of electrolytic charged specimens was reported in [26] for several similar ODS steel grades. Unfortunately, the microstructure was not characterized in detail. The influence of hydrogen upon sustained load cracking of a powder metallurgical (PM) bcc steel with fine dispersed Cr2O3 precipitations was investigated in [27]. In both investigations [26,27] the critical hydrogen content below which no failure occurred, was 10 times higher in the ODS materials compared to conventional non-ODS reference materials which was attributed to effective H-trapping especially at the oxide particles. The aim of this study was to investigate if a high amount of irreversible traps can effectively reduce the susceptibility to HEE by testing a ferritic ODS steel with nano-sized Y2O3 particles in high pressure hydrogen atmosphere.
2.
Experimental
The nominal and the measured chemical compositions of the material are given in Table 2. The steel bar was vacuum heat treated at 1100 C for 30 min followed by 750 C for 2 h and subsequent vacuum cooling.
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Transmission electron microscope (TEM) data and scanning TEM (STEM) data were collected on a spherical aberration corrected (Cs) JEOL 2100F TEM/STEM operating at 200 kV. The microscope is equipped with a Schottky field emission gun, a CEOS GmbH hexapole aberration corrector, high angle annular dark field (HAADF) detector, an Oxford energy dispersive x-ray spectrometer (EDS) and a Gatan Tridiem electron energy filter. The TEM/STEM samples were prepared by focused ion beam (FIB) thinning using a Carl Zeiss NVision CrossBeam FIB-SEM. HEE was evaluated by slow displacement rate tensile tests (longitudinal direction, 0.1 mm/min according ASTM G 142) in H2 gas (99.9999%) at 20 C and 10 MPa using cylindrical specimens (gauge length 30 mm, gauge diameter 5 mm). Reference tests were performed in air at ambient laboratory conditions. Tests in H2 were performed at the Materials Testing Institute University of Stuttgart (Germany) with a test apparatus described in [28]. Tests in air were performed at Opel central laboratories, Ruesselsheim, Germany. Fractographic analysis was performed by conventional SEM techniques.
3.
Results and discussion
3.1.
Microstructure characterization
Micrographs taken of the steel investigated are displayed in Fig. 1. As visible in the light micrograph in Fig. 1a, the microstructure has many precipitates distributed on the 2e8 mm length scale. The ferritic grain structure is visible in the SEM image (Fig. 1b) but few precipitates are discernible. The ferritic grains size ranges from 0.2 to 2 mm. From a comparison of the two micrographs, it is reasonable to assume that the region between the precipitates consists of several ferrite grains. It shall be noted here that no cavities were found in the microstructure indicating a density of nearly 100% of the powder metallurgical fabricated material. To identify the different precipitates, EDS elemental maps were collected and are displayed in Fig. 2. The precipitates appear streaked due to specimen drift encountered during the long data collection time. The maps reveal three types of precipitates. The most common precipitates are chromium based containing little other alloying elements. The second type of precipitate contains vanadium and chromium and the
Fig. 5 e EFTEM analysis of large carbide precipitates. a) bright field image, b) CrL map, c) CK map.
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Fig. 6 e EDS analysis of fine dispersed yttria precipitates. a) HAADF STEM image, the large black feature is a CreMn oxide comparable to the one shown in Fig. 3, b) Y map, c) O map.
third type of precipitates contains manganese, vanadium and chromium. The manganese containing precipitates are the least common but largest precipitates. The black appearing precipitates seen in Fig. 1b were all found to contain vanadium. A TEM micrograph and energy filtered elemental maps of the manganese containing precipitates are shown in Fig. 3. The bright field image (Fig. 3a) shows a high density of dislocations (dark) in the ferritic matrix. The manganese precipitates were only observed on the grain boundaries and identified as (Cr,Mn)xOy consistent to that previously reported [19]. The size of the (Cr,Mn)xOy precipitates range in the order of 10e500 nm. From the EDS analyses, the composition of the large (Cr,Mn)xOy precipitate shown in Fig. 3 was calculated as (Fe0.35Cr0.44Mn0.21)3.2O4. The energy filtered TEM (EFTEM) images (Fig. 3bef) revealed that large oxides have a uniform composition and that VN decorates the perimeter of such oxides. Selected area electron diffraction data (not shown) identified the oxide as Me3O4 (Me ¼ Cr, Fe or Mn) having a spinel structure ðFd 3mÞ and a lattice parameter of
˚ . High resolution TEM (Fig. 4) indicates that these a ¼ 8.4 A precipitates are incoherent to the ferritic matrix structure. Diffraction data (not shown) identified the nitride as VN having a rock salt structure ðFm 3mÞ with a lattice param˚ . The spinel and rock salt phases are coherent eter of a ¼ 4.1 A with a cube-cube orientation relationship. Several non-oxide precipitates can be found in the TEM images in Fig. 5. As in Fig. 3, the ferrite matrix has a high dislocation density. The majority of the non-oxide precipitates are Cr23C6 with a smaller number of VN precipitates. These precipitates are about around 100 nm in size. The Cr23C6 and VN precipitates are located on the grain boundaries, discontinuous discrete particles and incoherent to the ferritic matrix structure. The dispersed oxide precipitates, YxOy, are shown in Figs. 6 and 7. The dispersed oxides precipitates are mainly spherical particles finely dispersed in the ferritic matrix. While some regions were fairly devoid of the YxOy precipitates, other areas such as that shown in Fig. 7. have a large number of these precipitates, thus they are not uniformly distributed
Fig. 7 e STEM images of the fine oxide precipitates. a) HAADF STEM image, b) enlarged STEM HAADF image of a single fine oxide precipitate c) Fe EELS map, d) Cr EELS map, e) Mn EELS map, f) V EELS map, g) O EELS map, f) N EELS map.
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(Cr,Mn)xOy incoherent precipitate (diameter 10-500 nm) with VN attached.
Matrix (a ˜ 3 Å)
Grain boundary
Y precipitate (diameter 5-20 nm)
YxOy precipitate (diameter 5-20 nm)
CrxCy precipitate (diameter ~100 nm)
Fig. 8 e Schematic view of the detected precipitates.
throughout the ferritic matrix. It is suspected that the formation of larger grain boundary oxides, Me3O4, result in the reduction in number of the yttrium oxide precipitates in the ferritic matrix. The precipitates YxOy are about 5e20 nm in size and in areas with a large number of precipitates they have a spatial distribution in the order of 100 nm. As demonstrated by the electron energy loss spectroscopy (EELS) elemental maps in Fig. 6, the YxOy precipitates have a core shell structure. The shell, approximately 2 nm thick, is rich in Cr and V and the core has a slightly higher concentration of manganese, similar to those previously reported [23]. A schematic view of the microstructure is given in Fig. 8.
3.2.
detailed view of the H-assisted crack morphology is shown in Fig. 10d and e. This fracture surface is characterized by very small facets and a regular crack pattern. The size of the facets and of the crack pattern is on the order of 1e8 mm, which corresponds quite well with the separation between the grain boundary precipitates indicating a high fracture portion occurring along the grain boundaries decorated with the larger precipitates (Me3O4, VN and Cr23C6). On the other hand, areas with clear indications of plasticity are visible on the fracture surface (Fig. 10e) indicating transgranular fracture. Thus, it can be assumed that the overall fracture morphology is characterized by a mixture of intergranular H-assisted fracture and transgranular HELP fracture.
Tensile tests
The results of the tensile tests are summarized in Table 3. Yield (0.2YS) and ultimate tensile strength (UTS) were unaffected by the H2 atmosphere whereas elongation at fracture (El) and reduction in area (RA) were markedly reduced. The corresponding stressestrain diagrams are shown in Fig. 9. In air, uniform elongation is quite short (w4%) followed by a larger elongation of about 8% during necking which seems to be an effect of the low strain rate applied here [29]. In H2 atmosphere, uniform plastic strain is quite short as well. Since YS was unaffected by hydrogen, it can be assumed that H-assisted crack growth started at the onset of plastic straining. The corresponding fractographs are shown in Fig. 10. When tested in air, the surface is characterized by a cuttershaped fracture (Fig. 10a) and very fine dimples (Fig. 10b). In H2 atmosphere, a main H-assisted crack (the maximum depth of about 2 mm was approximated from Fig. 10c) propagated from the surface until the final cross section failed catastrophically due to overload (Fig. 10c). With this data and with the specimen dimensions and test data given in paragraph 2, the crack growth rate can be estimated to about 4 mm/s which is about 100 times faster compared to SUS 316L stainless steel [30]. A
3.3.
Role of irreversible traps
The following argumentation is based on a trap model which was extensively developed in [31] and summarized in [32]. Furthermore, it is assumed that HELP plays a dominant role in the deformation process, i.e. hydrogen is transported by moving dislocations. Furthermore, it is known from experiments performed in gaseous hydrogen atmosphere that H-assisted cracking starts from the surface. At long time exposure, e.g. automotive gaseous H2 applications, there is an unlimited source of H in combination with mechanical strains. This situation together with some main
Table 3 e Tensile test results in H2 gas (10 MPa) and ambient air at room temperature. Atmosphere
H2 Air
0.2YS
UTS
El
RA
MPa
MPa
%
%
963 950
1064 1070
4 13
10 63
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1200 1100 air, 0.1 MPa, 20°C
Engineering stress [MPa]
1000 900 800 700 H2, 10 MPa, 20°C
600 500 400 300 200 100 0 0
1
2
3
4
5
6
7
8
9
10
11
12
13
14
15
Engineering strain [%]
Fig. 9 e Engineering stressestrain diagram. The strain was calculated from the cross head displacement data.
dimensions is illustrated in Fig. 11 for the steel investigated here. Once the H2 molecule has dissociated on the strained surface, there are three main reaction paths. Path 1 is quite straight forward when the dissociation takes place near or at
an irreversible trap, e.g. an oxide particle. In this case, H is trapped right away without any influence on dislocation slip and the total amount of interstitial H is reduced. If this irreversible trap (here: mainly (Cr,Mn)xOy) is located at a grain boundary hydrogen is accumulated at the grain boundary enhancing the tendency for intergranular fracture. For paths 2 and 3, dissociation takes place somewhere in between irreversible traps. In these cases, H can be trapped at a moving dislocation and influence slip (HELP mechanism). When a moving dislocation meets a grain boundary precipitates, the dislocation most likely will lose the H proton (Hþ) to the irreversible trap further enhancing the tendency for intergranular fracture (path 2). On the other hand, when a moving dislocation meets a dispersed oxide particle (here YxOy), the dislocation will also lose the Hþ to the irreversible trap reducing the amount of transportable hydrogen (path 3). With the dimensions estimated by the microstructure analysis, dislocations can move in the order of 100e200 lattice sites (100 nm/2/ 0.3 nm ¼ 160) before the Hþ is transferred to either irreversible trap ((Cr,Mn)xOy or YxOy). It can be guessed that this distance is large enough for an H proton to significantly influence slip which is supported by the transgranular fracture portions described in paragraph 3.2.
Fig. 10 e (a) Overview of fracture surface tested in He, (b) detail of the fracture surface tested in He, (c) overview of the fracture surface tested in H2, area 1 denotes the primary H-assisted crack, (d) detail of area 1 of fracture surface tested in H2, (e) enlargement from (d).
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H H
1 ~ 0.1 nm
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precipitates which enhance the tendency for intergranular fracture.
2
H
H Surface
H G
2
Acknowledgements
3
H
H
Latticewith dislocation
F
The authors would like to thank Rainer Lindau from Forschungszentrum Karlsruhe, Germany for providing the test material and Nicholas Irish from GM R&D for providing the chemical analysis data. Strain
references
Fig. 11 e Schematic of hydrogen uptake and transport in a steel with grain boundary as well as fine dispersed irreversible traps and an unlimited hydrogen source. G, grain boundary precipitate. F, fine dispersed oxide, here YxOy.
Based on these results, it can be concluded that high energy trap sites can have positive and negative effects on the HEE resistance of steels. If a sufficient number of high energy trap sites are located at grain boundaries, the risk of intergranular fracture is high. On the other hand, if high energy trap sites are finely dispersed in the lattice, as transgranular HELP-assisted fracture morphology implies, then dislocation slip is indeed influenced by H but the total amount of interstitial H might be reduced below a threshold value resulting in unaffected stressestrain properties under the given test conditions. This argumentation was used previously to explain the unaffected stressestrain data of a precipitation hardened ferritic-pearlitic steel being tested in 10 MPa H2 atmosphere compared to laboratory air [33].
4.
Summary
The results found in this study can be summarized as follows: The microstructure of the ODS RAF steel studied in this work consists of a fine grain ferritic matrix, numerous grain boundary precipitates and dispersed yttrium oxide nano precipitates in the ferrite matrix. The grain boundary precipitates are predominately Me3O4, VN and Cr23C6. YS and UTS were unaffected by the H2 atmosphere. In H2, elongation at fracture reduced from 13% to 4% and reduction of area reduced from 63% to 10% as compared to the results obtained in air. In H2 atmosphere, the fracture morphology was found to be a mixture of intergranular H-assisted fracture and a smaller amount of transgranular HELP fracture. The sensitivity of the ODS RAF steel to hydrogen embrittlement is attributed to the large number grain boundary
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[17] Kimura A. Current status of reduced-activation ferritic/ martensitic steels R&D for fusion energy. Mater Trans 2005; 46(3):394e404. [18] El-Genk M, Tounier J-M. A review of refractory metal alloys and mechanically alloyed-oxide dispersion strengthened steels for space nuclear systems. J Nucl Mater 2005;340:93e112. [19] Lindau R, Mo¨slang A, Schirra M, Schlosmacher P, Klimenkov M. Mechanical and microstuctural properties of a hipped RAFM ODS-steel. J Nucl Mater 2002;307e311:769e72. [20] Capdevila C, Bhadeshia KDH. Manufacturing and microstructural evolution of mechanically alloyed oxide dispersion strengthened superalloys. Adv Eng Mater 2001;3 (9):647e56. [21] Eiselt ChCh, Klimenkov M, Lindau R, Mo¨slang A. Characteristic results and prospects of the 13Cr-1W-0.3Ti-0. 3Y2O3 ODS steel. J Nucl Mater 2009;386e388:525e8. [22] Eiselt ChCh, Klimenkov M, Lindau R, Mo¨slang A, Sandim HRZ, Padilha AF, et al. High-resolution transmission electron microscopy and electron backscatter diffraction in nanoscaled ferritic and ferritic-martensitic oxide dispersion strengthened-steels. J Nucl Mater 2009;385(2):231e5. [23] Klimenkov M, Lindau R, Mo¨slang A. New insights into the structure of ODS particles in the ODS-Eurofer alloy. J Nucl Mater 2009;386e388:553e6. [24] Klimenkov M, Lindau R, Mo¨slang A. TEM study of internal oxidation in an ODS-Eurofer alloy. J Nucl Mater 2009; 386e388:557e60.
[25] Esteban GA, Pen˜a A, Legarda F, Lindau R. Hydrogen transport and trapping in ODS-EUROFER. Fusion Eng Des 2007;82 (15e24):2634e40. [26] Lee JS, Kimura A, Ukai S, Fujukawa M. Effects of hydrogen on the mechanical properties of dispersion strengthened steels. J Nucl Mater 2004;329e333:1122e6. [27] Kimura Y, Sakai Y, Hara T, Belyakov A, Tsuzaki T. Hydrogen induced delayed fracture of ultrafine grained 0.6% O steel with dispersed oxide particles. Scripta Mater 2003;49: 1111e6. [28] Deimel P, Hanisch C. Tests on the steels 15 MnNi6 3 and X56TM in high pressure hydrogen gas of high purity. Int J Hydrogen Energy 1989;14:147e51. [29] Sokolov MA, Hoelzer DT, Stoller RE, McClintock DA. Fracture toughness and tensile properties of nano-structured ferritic steel 12YWT. J Nucl Mater 2007;367e370:213e6. [30] Michler T, Lee Y, Gangloff RP, Naumann J. Influence of macro segregation on hydrogen environment embrittlement of SUS 316L stainless steel. Int J Hydrogen Energy 2009;34:3201e9. [31] Pressouyre GM. The role of trapping on hydrogen transport and embrittlement. PhD thesis Carnegie-Mellon University, 1977. [32] Pressouyre GM. Trap theory of hydrogen embrittlement. Acta Metall 1980;28:859e911. [33] Michler T, Naumann J. Microstructural aspects upon hydrogen environment embrittlement of various bcc steels. Int J Hydrogen Energy 2010;35(2):821e32.
Available at www.sciencedirect.com
journal homepage: www.elsevier.com/locate/he
Hydrogen environment embrittlement of an ODS RAF steel e Role of irreversible hydrogen trap sites Thorsten Michler a,*, Michael P. Balogh b a b
Adam Opel GmbH, Ru¨sselsheim, Germany General Motors Research and Development, Warren, MI, USA
article info
abstract
Article history:
The microstructure and the effects of 10 MPa hydrogen atmosphere on the tensile prop-
Received 3 May 2010
erties of a oxide dispersion strengthened (ODS) reduced activation ferritic (RAF) steel were
Received in revised form
investigated. The microstructure consists of a fine grained ferritic matrix with Me3O4
24 June 2010
(Me ¼ Cr, Fe or Mn), VN and Cr23C6 grain boundary precipitates as well as dispersed yttrium
Accepted 24 June 2010
oxide nano precipitates in the ferritic matrix. The yield and ultimate tensile strength were
Available online 1 August 2010
unaffected by the H2 atmosphere whereas elongation at fracture and reduction in area were markedly reduced. In H2 atmosphere, the fracture morphology was found to be
Keywords:
a mixture of intergranular H-assisted fracture and a smaller amount of transgranular
Hydrogen
hydrogen enhanced localized plasticity (HELP) fracture. The sensitivity of the ODS RAF steel
Hydrogen traps
to hydrogen embrittlement is attributed to the large number grain boundary precipitates
Hydrogen embrittlement
which enhance the tendency for intergranular fracture.
ODS steel
ª 2010 Professor T. Nejat Veziroglu. Published by Elsevier Ltd. All rights reserved.
Slow displacement rate tensile test
1.
Introduction
Hydrogen environment embrittlement (HEE) is a well known phenomenon 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 20e875 bar. Common materials for the design of hydrogen wetted components are CreNi austenitic stainless steels (SS). For stationary hydrogen tanks where cost and weight are of minor importance, 18Cr-10Ni-2Mo steels like DIN 1.4404 (AISI 316L) or DIN 1.4571 (AISI 316 Ti) are widely and successfully used. The best resistance to HEE under severe conditions is reached for Ni contents higher than 12.5 wt% in solution treated condition [1,2]. Nickel and molybdenum are the cost drivers in SS which makes these grades unattractive for automotive mass production. Furthermore, the strength of SS is quite low (yield strength from about 200 to 250 MPa)
resulting in thick tank walls and ultimately making such tanks not only expensive but also quite heavy. Significant cost and weight savings would be possible if a higher strength steel could be utilized. Unfortunately, for conventional steels, hydrogen enhances the susceptibility of high strength steels to brittle fracture, i.e. the higher the strength the higher the susceptibility [3] in addition to the fact that high strength steels are intrinsically more prone to brittle cracking than lower strength steels. Today, the most accepted explanation for HEE is the hydrogen enhanced localized plasticity (HELP) model. According to HELP, dissolved H enhances dislocation mobility and slip planarity [4e7], which can lead to heterogeneously localized plastic strain and stress concentrations presumably sufficient to enable subcritical crack growth. Due to enhanced-heterogeneous planar slip with H present, the macroscopic fracture appearance is “brittle” and involves
* Corresponding author. E-mail address: [email protected] (T. Michler). 0360-3199/$ e see front matter ª 2010 Professor T. Nejat Veziroglu. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.ijhydene.2010.06.071
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Table 1 e Some reversible and irreversible traps and their binding energies relative to atomic hydrogen. Trap
Binding energy Reference [kJ/mol]
C interstitial Interstitial site in iron based alloy Dislocation in Fe Screw dislocation core Grain boundary Fe oxide interface Y2O3 interface Al2O3 interface Cr carbide interface Fe3C interface
3 4e8 24e26 20e30 18e53 47 70 79 67 84
[12] [12] [11,12] [11] [11] [11] [11] [11] [13] [11]
either slip band interface cracking or microcrack formation at slip band, twin and/or grain boundary intersections [8e10]. Another characteristic of hydrogen in steels is that hydrogen is not distributed homogeneously in the lattice. In general, hydrogen is trapped at higher binding energy
Table 2 e Chemical composition.
C Si Mn Cr W V Ta Y 2O 3
Nom. composition, wt %
Meas. composition, wt %
0.1 0.1 0.4 9 1.1 0.2 0.15 0.3e0.5
Not measured 0.05 0.38 8.7 0.66 0.19 0.04 0.25 (0.22 Y)
sites in comparison to normal interstitial sites. Trapping sites fall into two classes depending on the binding energy. Traps with binding energies less than w50 kJ/mol are called reversible traps, whereas those with higher energies are termed irreversible traps. Reversible traps are characterized by the ability to capture and release hydrogen atoms. Due to their high binding energies, irreversible traps only
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capture hydrogen until they become saturated and release of hydrogen atoms is very difficult if not impossible. As shown in Table 1, low binding energies are reported for common structural features like grain boundaries, dislocations or substitutional elements while high binding energies are found for vacancies and precipitates like carbides and oxides [11e13]. Recent studies on carbidic precipitates have shown that the trapping capacity also depends on the lattice mismatch, i.e. whether the carbides are coherent or incoherent [14,15]. The role of trapping upon hydrogen embrittlement is not clear yet. Trapping might have a positive effect since the amount of diffusible hydrogen is strongly reduced. On the other hand, trapping might be detrimental when hydrogen weakens the structural integrity, e.g. at grain boundaries. The concept of incorporating irreversible H traps into the lattice to reduce HEE was proposed quite some time ago (see review [16] and references therein) but systematic investigations are missing. To reduce the amount of diffusible hydrogen in the lattice, traps need to be homogenously distributed, have a high binding energy, available in sufficient quantity and separated by a minimal distance [16]. IF HELP is assumed to be the dominant mechanism, the effect of H on dislocation mobility might be reduced when H is trapped in irreversible traps. Since the mobility of H in HEE experiments at room temperature is only a few mm, fine dispersed nano precipitates could be promising. Reduced activation ferritic/martensitic (RAFM) steels are materials for structural applications in future fusion reactors. The main requirements are a high resistance to both transmutation He-induced embrittlement as well as to neutron irradiation embrittlement which decreases the fracture toughness of many materials and causes an increase in ductile-to-brittle transition temperature. One candidate material is EUROFER 97 with a base composition of 0.1 wt% C, 9 wt% Cr, 0.4 wt% Mn and 0.1 wt% Si. Other (micro) alloying elements are 1.1 wt% W, 0.2 wt% V and 0.15 wt% Ta added to reduce neutron radiation induced swelling. About 9 wt% Cr is added to reduce irradiation hardening. Furthermore, the superior performance of RAFM steels is attributed to the high helium trapping capacity of the structure, which contains martensite laths, precipitates and carbides as traps with high binding energies [17,18]. To increase the operating temperature and thus the efficiency of the reactors, mechanical
Fig. 1 e Macroscopic structure of the ODS EUROFER steel. a) Light micrograph V2A etchant, b). SEM image, the vertical structures are artifacts from the FIB polishing process.
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Fig. 2 e SEM EDS elemental maps collected from the FIB polished specimen, see Fig. 1b. a) Fe map, b) Cr map, c) Mn map, d) V map.
milling is used to embed finely dispersed oxide particles (e.g. Y2O3) into the base alloy (e.g. EUROFER 97) which increases the creep strength [19,20]. Over the past years, oxide dispersion strengthened (ODS) alloys have attracted much attention and
their properties have been characterized in detail [17e19,21e24]. Due to Y2O3 alloying, the structure is entirely ferritic [19] and thus, such steels are denoted reduced activation ferritic (RAF) steels.
Fig. 3 e EFTEM images of a large oxide precipitate. a) bright field image, b) CrL map, c) MnL map, d) VL map, e) OK map, f) NK map.
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Fig. 4 e High resolution TEM image of a large oxide precipitate.
Hydrogen trapping and transport in ODS-EUROFER was investigated in [25] showing that Y2O3 precipitates act as irreversible hydrogen traps. Severe influence of hydrogen upon tensile ductility of electrolytic charged specimens was reported in [26] for several similar ODS steel grades. Unfortunately, the microstructure was not characterized in detail. The influence of hydrogen upon sustained load cracking of a powder metallurgical (PM) bcc steel with fine dispersed Cr2O3 precipitations was investigated in [27]. In both investigations [26,27] the critical hydrogen content below which no failure occurred, was 10 times higher in the ODS materials compared to conventional non-ODS reference materials which was attributed to effective H-trapping especially at the oxide particles. The aim of this study was to investigate if a high amount of irreversible traps can effectively reduce the susceptibility to HEE by testing a ferritic ODS steel with nano-sized Y2O3 particles in high pressure hydrogen atmosphere.
2.
Experimental
The nominal and the measured chemical compositions of the material are given in Table 2. The steel bar was vacuum heat treated at 1100 C for 30 min followed by 750 C for 2 h and subsequent vacuum cooling.
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Transmission electron microscope (TEM) data and scanning TEM (STEM) data were collected on a spherical aberration corrected (Cs) JEOL 2100F TEM/STEM operating at 200 kV. The microscope is equipped with a Schottky field emission gun, a CEOS GmbH hexapole aberration corrector, high angle annular dark field (HAADF) detector, an Oxford energy dispersive x-ray spectrometer (EDS) and a Gatan Tridiem electron energy filter. The TEM/STEM samples were prepared by focused ion beam (FIB) thinning using a Carl Zeiss NVision CrossBeam FIB-SEM. HEE was evaluated by slow displacement rate tensile tests (longitudinal direction, 0.1 mm/min according ASTM G 142) in H2 gas (99.9999%) at 20 C and 10 MPa using cylindrical specimens (gauge length 30 mm, gauge diameter 5 mm). Reference tests were performed in air at ambient laboratory conditions. Tests in H2 were performed at the Materials Testing Institute University of Stuttgart (Germany) with a test apparatus described in [28]. Tests in air were performed at Opel central laboratories, Ruesselsheim, Germany. Fractographic analysis was performed by conventional SEM techniques.
3.
Results and discussion
3.1.
Microstructure characterization
Micrographs taken of the steel investigated are displayed in Fig. 1. As visible in the light micrograph in Fig. 1a, the microstructure has many precipitates distributed on the 2e8 mm length scale. The ferritic grain structure is visible in the SEM image (Fig. 1b) but few precipitates are discernible. The ferritic grains size ranges from 0.2 to 2 mm. From a comparison of the two micrographs, it is reasonable to assume that the region between the precipitates consists of several ferrite grains. It shall be noted here that no cavities were found in the microstructure indicating a density of nearly 100% of the powder metallurgical fabricated material. To identify the different precipitates, EDS elemental maps were collected and are displayed in Fig. 2. The precipitates appear streaked due to specimen drift encountered during the long data collection time. The maps reveal three types of precipitates. The most common precipitates are chromium based containing little other alloying elements. The second type of precipitate contains vanadium and chromium and the
Fig. 5 e EFTEM analysis of large carbide precipitates. a) bright field image, b) CrL map, c) CK map.
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Fig. 6 e EDS analysis of fine dispersed yttria precipitates. a) HAADF STEM image, the large black feature is a CreMn oxide comparable to the one shown in Fig. 3, b) Y map, c) O map.
third type of precipitates contains manganese, vanadium and chromium. The manganese containing precipitates are the least common but largest precipitates. The black appearing precipitates seen in Fig. 1b were all found to contain vanadium. A TEM micrograph and energy filtered elemental maps of the manganese containing precipitates are shown in Fig. 3. The bright field image (Fig. 3a) shows a high density of dislocations (dark) in the ferritic matrix. The manganese precipitates were only observed on the grain boundaries and identified as (Cr,Mn)xOy consistent to that previously reported [19]. The size of the (Cr,Mn)xOy precipitates range in the order of 10e500 nm. From the EDS analyses, the composition of the large (Cr,Mn)xOy precipitate shown in Fig. 3 was calculated as (Fe0.35Cr0.44Mn0.21)3.2O4. The energy filtered TEM (EFTEM) images (Fig. 3bef) revealed that large oxides have a uniform composition and that VN decorates the perimeter of such oxides. Selected area electron diffraction data (not shown) identified the oxide as Me3O4 (Me ¼ Cr, Fe or Mn) having a spinel structure ðFd 3mÞ and a lattice parameter of
˚ . High resolution TEM (Fig. 4) indicates that these a ¼ 8.4 A precipitates are incoherent to the ferritic matrix structure. Diffraction data (not shown) identified the nitride as VN having a rock salt structure ðFm 3mÞ with a lattice param˚ . The spinel and rock salt phases are coherent eter of a ¼ 4.1 A with a cube-cube orientation relationship. Several non-oxide precipitates can be found in the TEM images in Fig. 5. As in Fig. 3, the ferrite matrix has a high dislocation density. The majority of the non-oxide precipitates are Cr23C6 with a smaller number of VN precipitates. These precipitates are about around 100 nm in size. The Cr23C6 and VN precipitates are located on the grain boundaries, discontinuous discrete particles and incoherent to the ferritic matrix structure. The dispersed oxide precipitates, YxOy, are shown in Figs. 6 and 7. The dispersed oxides precipitates are mainly spherical particles finely dispersed in the ferritic matrix. While some regions were fairly devoid of the YxOy precipitates, other areas such as that shown in Fig. 7. have a large number of these precipitates, thus they are not uniformly distributed
Fig. 7 e STEM images of the fine oxide precipitates. a) HAADF STEM image, b) enlarged STEM HAADF image of a single fine oxide precipitate c) Fe EELS map, d) Cr EELS map, e) Mn EELS map, f) V EELS map, g) O EELS map, f) N EELS map.
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(Cr,Mn)xOy incoherent precipitate (diameter 10-500 nm) with VN attached.
Matrix (a ˜ 3 Å)
Grain boundary
Y precipitate (diameter 5-20 nm)
YxOy precipitate (diameter 5-20 nm)
CrxCy precipitate (diameter ~100 nm)
Fig. 8 e Schematic view of the detected precipitates.
throughout the ferritic matrix. It is suspected that the formation of larger grain boundary oxides, Me3O4, result in the reduction in number of the yttrium oxide precipitates in the ferritic matrix. The precipitates YxOy are about 5e20 nm in size and in areas with a large number of precipitates they have a spatial distribution in the order of 100 nm. As demonstrated by the electron energy loss spectroscopy (EELS) elemental maps in Fig. 6, the YxOy precipitates have a core shell structure. The shell, approximately 2 nm thick, is rich in Cr and V and the core has a slightly higher concentration of manganese, similar to those previously reported [23]. A schematic view of the microstructure is given in Fig. 8.
3.2.
detailed view of the H-assisted crack morphology is shown in Fig. 10d and e. This fracture surface is characterized by very small facets and a regular crack pattern. The size of the facets and of the crack pattern is on the order of 1e8 mm, which corresponds quite well with the separation between the grain boundary precipitates indicating a high fracture portion occurring along the grain boundaries decorated with the larger precipitates (Me3O4, VN and Cr23C6). On the other hand, areas with clear indications of plasticity are visible on the fracture surface (Fig. 10e) indicating transgranular fracture. Thus, it can be assumed that the overall fracture morphology is characterized by a mixture of intergranular H-assisted fracture and transgranular HELP fracture.
Tensile tests
The results of the tensile tests are summarized in Table 3. Yield (0.2YS) and ultimate tensile strength (UTS) were unaffected by the H2 atmosphere whereas elongation at fracture (El) and reduction in area (RA) were markedly reduced. The corresponding stressestrain diagrams are shown in Fig. 9. In air, uniform elongation is quite short (w4%) followed by a larger elongation of about 8% during necking which seems to be an effect of the low strain rate applied here [29]. In H2 atmosphere, uniform plastic strain is quite short as well. Since YS was unaffected by hydrogen, it can be assumed that H-assisted crack growth started at the onset of plastic straining. The corresponding fractographs are shown in Fig. 10. When tested in air, the surface is characterized by a cuttershaped fracture (Fig. 10a) and very fine dimples (Fig. 10b). In H2 atmosphere, a main H-assisted crack (the maximum depth of about 2 mm was approximated from Fig. 10c) propagated from the surface until the final cross section failed catastrophically due to overload (Fig. 10c). With this data and with the specimen dimensions and test data given in paragraph 2, the crack growth rate can be estimated to about 4 mm/s which is about 100 times faster compared to SUS 316L stainless steel [30]. A
3.3.
Role of irreversible traps
The following argumentation is based on a trap model which was extensively developed in [31] and summarized in [32]. Furthermore, it is assumed that HELP plays a dominant role in the deformation process, i.e. hydrogen is transported by moving dislocations. Furthermore, it is known from experiments performed in gaseous hydrogen atmosphere that H-assisted cracking starts from the surface. At long time exposure, e.g. automotive gaseous H2 applications, there is an unlimited source of H in combination with mechanical strains. This situation together with some main
Table 3 e Tensile test results in H2 gas (10 MPa) and ambient air at room temperature. Atmosphere
H2 Air
0.2YS
UTS
El
RA
MPa
MPa
%
%
963 950
1064 1070
4 13
10 63
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1200 1100 air, 0.1 MPa, 20°C
Engineering stress [MPa]
1000 900 800 700 H2, 10 MPa, 20°C
600 500 400 300 200 100 0 0
1
2
3
4
5
6
7
8
9
10
11
12
13
14
15
Engineering strain [%]
Fig. 9 e Engineering stressestrain diagram. The strain was calculated from the cross head displacement data.
dimensions is illustrated in Fig. 11 for the steel investigated here. Once the H2 molecule has dissociated on the strained surface, there are three main reaction paths. Path 1 is quite straight forward when the dissociation takes place near or at
an irreversible trap, e.g. an oxide particle. In this case, H is trapped right away without any influence on dislocation slip and the total amount of interstitial H is reduced. If this irreversible trap (here: mainly (Cr,Mn)xOy) is located at a grain boundary hydrogen is accumulated at the grain boundary enhancing the tendency for intergranular fracture. For paths 2 and 3, dissociation takes place somewhere in between irreversible traps. In these cases, H can be trapped at a moving dislocation and influence slip (HELP mechanism). When a moving dislocation meets a grain boundary precipitates, the dislocation most likely will lose the H proton (Hþ) to the irreversible trap further enhancing the tendency for intergranular fracture (path 2). On the other hand, when a moving dislocation meets a dispersed oxide particle (here YxOy), the dislocation will also lose the Hþ to the irreversible trap reducing the amount of transportable hydrogen (path 3). With the dimensions estimated by the microstructure analysis, dislocations can move in the order of 100e200 lattice sites (100 nm/2/ 0.3 nm ¼ 160) before the Hþ is transferred to either irreversible trap ((Cr,Mn)xOy or YxOy). It can be guessed that this distance is large enough for an H proton to significantly influence slip which is supported by the transgranular fracture portions described in paragraph 3.2.
Fig. 10 e (a) Overview of fracture surface tested in He, (b) detail of the fracture surface tested in He, (c) overview of the fracture surface tested in H2, area 1 denotes the primary H-assisted crack, (d) detail of area 1 of fracture surface tested in H2, (e) enlargement from (d).
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H H
1 ~ 0.1 nm
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precipitates which enhance the tendency for intergranular fracture.
2
H
H Surface
H G
2
Acknowledgements
3
H
H
Latticewith dislocation
F
The authors would like to thank Rainer Lindau from Forschungszentrum Karlsruhe, Germany for providing the test material and Nicholas Irish from GM R&D for providing the chemical analysis data. Strain
references
Fig. 11 e Schematic of hydrogen uptake and transport in a steel with grain boundary as well as fine dispersed irreversible traps and an unlimited hydrogen source. G, grain boundary precipitate. F, fine dispersed oxide, here YxOy.
Based on these results, it can be concluded that high energy trap sites can have positive and negative effects on the HEE resistance of steels. If a sufficient number of high energy trap sites are located at grain boundaries, the risk of intergranular fracture is high. On the other hand, if high energy trap sites are finely dispersed in the lattice, as transgranular HELP-assisted fracture morphology implies, then dislocation slip is indeed influenced by H but the total amount of interstitial H might be reduced below a threshold value resulting in unaffected stressestrain properties under the given test conditions. This argumentation was used previously to explain the unaffected stressestrain data of a precipitation hardened ferritic-pearlitic steel being tested in 10 MPa H2 atmosphere compared to laboratory air [33].
4.
Summary
The results found in this study can be summarized as follows: The microstructure of the ODS RAF steel studied in this work consists of a fine grain ferritic matrix, numerous grain boundary precipitates and dispersed yttrium oxide nano precipitates in the ferrite matrix. The grain boundary precipitates are predominately Me3O4, VN and Cr23C6. YS and UTS were unaffected by the H2 atmosphere. In H2, elongation at fracture reduced from 13% to 4% and reduction of area reduced from 63% to 10% as compared to the results obtained in air. In H2 atmosphere, the fracture morphology was found to be a mixture of intergranular H-assisted fracture and a smaller amount of transgranular HELP fracture. The sensitivity of the ODS RAF steel to hydrogen embrittlement is attributed to the large number grain boundary
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