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The effect of hydrogen on the deterioration of austenitic steels during wear at cryogenic temperature
Wear 259 (2005) 424–431
The effect of hydrogen on the deterioration of austenitic steels during wear at cryogenic temperature H. Pinto a,∗ , A. Pyzalla a , R. B¨uscher b , A. Fischer b , K. Aßmus c , W. H¨ubner c a
Vienna University of Technology, TU Wien, Institute for Materials Science and Technology (E308), Karlsplatz 13, A-1040 Vienna, Austria b University Duisburg-Essen, Ipe-Werkstofftechnik II, Lotharstreet 1, 47057 Duisburg, Germany c Federal Institute for Materials Research and Testing (BAM), Unter den Eichen 44-46, 12200 Berlin, Germany Received 16 August 2004; received in revised form 3 February 2005; accepted 4 February 2005 Available online 10 May 2005
Abstract Hydrogen represents an important alternative to fossil fuels. Hydrogen storage is possible as a gas, at room temperature (RT) at about 20 MPa pressure, and in a liquefied form, at cryogenic temperatures of about 20 K. The latter form is particularly attractive due to the possibility of stocking a large quantity of hydrogen within a small volume. In moving parts (e.g. of transport vehicles) cryogenic temperature and the presence of hydrogen strongly enhance wear processes and subsequently component failure. The present work deals with the deformation behaviour and the microstructural deterioration of austenitic CrNi- and CrMn high nitrogensteels during friction in liquid hydrogen at 20 K. The modified microstructure within the wear scar is studied by scanning electron microscopy and X-ray diffraction methods. Diffraction studies of wear scars reveal the importance of twinning during deformation at 20 K. This increase of twinning can be attributed to a hydrogen-induced reduction of stacking fault energy (SFE) in the austenitic steels. Interactions between twin boundaries and planar dislocation structures along with locally increased stresses led to the formation of extensive crack networks. The amount of hydrogen-induced surface cracks depends on the alloy composition and is not necessarily correlated to the wear resistance of the austenitic steels. © 2005 Elsevier B.V. All rights reserved. Keywords: Austenitic steel; Stacking fault energy; Hydrogen; Wear; Cryotechnology
1. Introduction Austenitic steels are materials commonly used in engines, which operate at cryogenic temperatures. These steels combine a high mechanical strength with sufficient ductility at extremely low temperatures. Cryogenic temperatures are encountered, e.g. in aerospace technology, where liquid hydrogen (LH2 ) became the standard propellant. Recently many efforts are being done to employ liquid hydrogen in fuel cells. Since hydrogen represents an important alternative to fossil fuels in the automotive sector, the production of hydrogen on a large scale along with its compatibility to the current technology of internal combustion engine are of great technological concern. Hydrogen storage is possible in tanks as a ∗
Corresponding author. Tel.: +43 1 58801 30836; fax: +43 1 58801 30897. E-mail address: [email protected] (H. Pinto).
0043-1648/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.wear.2005.02.057
gas, at room temperature (RT) at about 20 MPa pressure, and in a liquefied form, at cryogenic temperatures of about 20 K. The latter form is particularly attractive due to the possibility of stocking of a large quantity of hydrogen within a small volume. Due to the impracticable use of lubricants at extremely low temperatures, moving machine components are exposed to high wear rates. Wear processes are often correlated to an unexpected failure of sealings, valves and bearings. This could lead to an escape of hydrogen and the formation of an explosive atmosphere. Furthermore austenitic steels may undergo a deformationinduced martensitic-phase transformation [1], which is especially favoured by low temperatures. The formation of martensite within the worn surface enhances strength but also causes embrittlement and thus influences the degradation process of cryogenic machine components. Hence, the
H. Pinto et al. / Wear 259 (2005) 424–431
lifetime of tribo-systems strongly depends on the microstructural changes within the near-surface region. The aim of the present work is the study of the influence of the chemical composition and the tribological stress conditions on the stability of the austenitic microstructure, the deterioration process and the wear behaviour during friction in LH2 of five different austenitic steels, which are currently in use. They provide a continuous variation of the stacking fault energy (SFE) as well as of the austenite stability mainly due to their different Ni-, Mn- and N-contents. The investigations attempt to correlate the SFE (and its variation with the adsorbed hydrogen content at distinct loading conditions) to the microstructural modifications and the wear performance.
2. Experimental Five different austenitic steels were chosen for the present investigations. Their chemical compositions are shown in Table 1. The SFE at room temperature were determined as described in [2]. The austenitic steels chosen differ in their Cr, Ni, Mn and N contents. Thus, they have strongly different stacking fault energies. According to [3], the susceptibility to deformationinduced martensitic transformation increases with decreasing SFE. In [4,5], it has been shown that due to its influence on the active deformation mechanisms the SFE plays an important role in the wear behaviour. The other microstructural parameters, such as grain size and volume fraction of inclusions and ␦-ferrite, which also affect the mechanical strength of austenitic steels, did not vary considerably for the different alloys investigated. In order to simulate wear processes at cryogenic temperatures, ball-on-disc sliding friction tests have been carried out using a special cryotribometer [6] in a liquid hydrogen LH2 (20 K) atmosphere at the Federal Institute for Materials Research and Testing (BAM) in Berlin, Germany. The loading parameters for the sliding tests were: normal force FN = 5 and 10 N, sliding speed, v = 0.06 and 0.2 m/s and sliding distance s = 1000 m. An aluminium oxide ceramic ball was employed as counterbody in order to minimize the transfer layer between the two sliding bodies. The worn surfaces were inspected using optical and scanning electron microscopy (SEM). The wear resistance was evaluated by means of the worn area on the ball specimen using SEM. Analysis of loss of weight and scar depth led to
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very high uncertainties due to the small scar dimensions. In order to determine the depth of the deformed near-surface layer, microhardness tests HV0.01 at the CrNi steel samples and nanohardness tests using a Berkovich identer at the CrMn steel samples were carried out. Nanohardness measurements were necessary due to extremely thin layers, which were formed on the surface in the steel 1.4452. For the nanohardness tests, a penetration depth of 0.3 m was chosen. Synchrotron X-ray diffraction studies were performed at the experimental station G3, at HASYLAB, DESY, in Hamburg, Germany, using a wavelength similar to Co K␣radiation. The strain-induced martensite formation within the wear scar, the deformation-induced lattice defects were quantified. For the evaluation of the martensite content after friction, the texture in each phase was taken into account [7,8]. The recorded diffractograms from the wear scar have been subjected to a Rietveld refinement procedure using the freeware material analysis using diffraction (MAUD) [9–11] in order to analyze the deformed microstructure. The instrument-broadening function has been determined by means of a standard silicon powder with a well-known grain size distribution. The SFE and the dislocation densities were evaluated from the microstrains and stacking fault probabilities determined by the Rietveld refinement procedure [4,12], as follows: K111 ω0 G(1 1 1) a0 A−0.37 ε250 1 1 1 SFE = (1) √ α π 3 where ε250 1 1 1 is the microstrain along the 1 1 1 crystallographic direction, α the stacking fault probability, K111 ω0 the proportionality constant, G(1 1 1) the shear modulus in the (1 1 1) fault plane = 1/3(c44 + c11 − c12 ), a0 the lattice constant, A the Zener anisotropy = 2c44 /(c11 − c12 ) and cij is the elastic stiffness coefficient. ρ=
48πα(SFE) (Ke + Ks )a02 d1 1 1
(2)
where ρ is the dislocation density, and d1 1 1 is the interplanar spacing and c 1/2 44 (c11 − c12 ) Ks = 2 and 1/2 c44 (c11 − c12 ) . Ke = (c11 + c12 ) c11 (c11 + c12 + 2c44 )
Table 1 Chemical composition of the investigated materials Materiala
SFERT (mJ/m2 )
Cr
Ni
Mn
Mo
C
N
Fe
1.4452 1.4301 1.4439 1.4876 1.4591
12 21 30 42 71
16.0–20.0 17.0–19.5 16.5–18.5 20.5 32.85
– 8.0–10.5 12.5–14.5 30.65 30.95
12.0–16.0 2.0 1.5 0.69 –
2.5–4.2 – 4.0–4.5 – 1.67
<0.15 <0.07 0.03 0.061 0.007
0.75–1.0 <0.11 0.12–0.22 – 0.39
Bal. Bal. Bal. Bal. Bal.
a
German steel grades.
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Fig. 1. Effect of increasing sliding velocities and normal forces on the wear resistance.
3. Results and discussion 3.1. Wear behaviour The wear performance during sliding friction tests in LH2 was evaluated by means of the wear scars created on the counterbody. These worn areas produced on the Al2 O3 -balls were chosen for the evaluation of the wear amount, since on the steel specimens, the scar depths are very small (comparable to the roughness of the polished surface) and the scar widths may lead to an erroneous interpretation due to the material transfer between disc and counterbody during adhesive wear. Fig. 1 displays the variation of the scar widths on the ball specimen with the tribological stress conditions for the different austenitic steels. The error bars give the standard deviation between the diameters of the scars measured along and perpendicular to the sliding direction. Results reveal that at mild tribological stressing (FN = 5 N, v = 0.06 m/s) the wear resistance of all investigated alloys is nearly the same. An increase of the sliding speed (FN = 5 N, v = 0.2 m/s) leads to a reduction of wear in all N-alloyed steels and to a linear increase of wear in the N-free 1.4876. On
the other hand, distinct wear performances for each investigated austenitic steel appear at more severe stress conditions (FN = 10 N, v = 0.2 m/s). The metastable CrNi steel 1.4301 and the CrMnN-steel 1.4452 with the lowest SFE present the best wear resistance The second best wear performance is found for the Mo-alloyed steel 1.4439 with about 0.17% N and moderate SFE (Fig. 1). For these three steel grades (1.4301, 1.4452 and 1.4439) the increase of normal force and sliding speed does not affect the material removal. A high SFE leads clearly to a loss of wear resistance. The highalloyed steel 1.4591 with 0.39% N exhibits an intermediate wear performance, whilst the Ni-free steel 1.4876 appears with the broadest scar. These distinct wear behaviours suggest an important role of the strain hardening by martensite formation on the wear performance of the metastable alloy 1.4301. In addition a positive effect of Mo- and N-alloying in CrNi steels can be observed. Moderate Mo- and N-contents are known to increase the yield strength, while sufficient toughness remains. This explains the enhanced wear resistance at cryogenic temperatures. A high N-content (1.0%) in the CrMn steel can result in loss of ductility [13]. However, the increased mechanical strength of the worn surface layer may be responsible for the best wear resistance of the CrMnN steel 1.4452. Fig. 2 shows typical images of the topography of the wear scars in inert environments. By an adhesive wear mechanism, particles are detached from the tribologically modified zone, then transferred to the counterbody and finally, re-transferred to the sample surface. The wear particles are generated as soon as the mechanical strength reaches its limit during the shear deformation of the surface. Crack formation could not be detected even at the lowest temperature, i.e. LHe (4.2 K), in all investigated alloys. In liquid hydrogen, the deformation behaviour in the contact zone seems to be different, as shown in Figs. 3–5. Crack formation is observed in particles deposited as well as in the worn matrix. The cracking density depends on the alloy and the loading conditions. In our previous works [14–17], we reported that an appreciable decrease of the SFE along with a very high yield strength and texture hardening favour cracking at the worn
Fig. 2. Wear scars after friction in LHe (FN = 5 N, v = 0.2 m/s): (a) steel 1.4301 and (b) steel 1.4439.
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Fig. 3. Steel 1.4452 after testing in LH2 at different loading conditions.
Fig. 4. Steel 1.4876 after testing in LH2 at different loading conditions.
Fig. 5. Steel 1.4591 after testing in LH2 at different loading conditions.
surface during sliding in LH2 by the pressure build-up mechanism [18,19]. In the present work, systematic SEM studies of the wear scars were carried out in order to evaluate the effect of the stress conditions, and the chemical composition on the cracking process for the different austenitic alloys. Figs. 3–5 display the worn surface of the austenitic steels 1.4452, 1.4876 and 1.4591 after friction at different conditions. In 1.4452, cracking occurs at all loading conditions and is intensified by increasing loadings. In turn, only very few cracks are formed in 1.4876 at the slightest applied stress condition of 5 N and 0.06 m/s. Heavier loading does not increase the crack formation but rather suppresses it. The steel 1.4591 at 5 N and 0.06 m/s is more sensitive to cracking than 1.4876 at 5 N and 0.06 m/s but similar to the later presents reduced cracking at heavier stress conditions. This suggests that in case of 1.4876 and 1.4591, the more pronounced deformation and severe loading conditions do not enhance the hydrogen uptake but likely facilitates the hydrogen release. It is assumed that hydrogen is continuously driven out and is not able to be soluted in the
austenitic crystal lattice. In all samples, the observed cracks formed within a nanocrystalline debris layer on the worn surface (Fig. 6). Therefore, crack appearance is not necessarily connected to a low wear resistance at cryogenic temperatures. As shown
Fig. 6. Deteriorated contact layer and crack network in the steel 1.4301 after friction in LH2 (FN = 5 N, v = 0.06 m/s).
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Fig. 7. Hardness–depth profiles of the near-surface zones after sliding in LH2 with FN = 5 N and v = 0.06 m/s.
in Fig. 1, 1.4876 and 1.4591 exhibit unsatisfactory wear performances but are less embrittled in presence of hydrogen. On the contrary, the metastable 1.4301, the steel 1.4439 and even the steel 1.4452 with the best wear performance reveal a very high crack density. The in-depth extension of the strengthened near-surface zone in the different austenitic alloys after friction in LH2 was evaluated by means of micro- and nanohardness tests. Fig. 7 shows the hardness–depth profiles for the steels 1.4452, 1.4301 and 1.4439 after tribological stressing with 5 N and 0.06 m/s. The depth of the hardened zone amounts to about 50 m for both CrNi steels (1.4301 and 1.4439), but the hardened zone is clearly thinner for the CrMnN steel 1.4452 (approximately 5 m). This reflects a pronounced increase of mechanical strength accompanied by a significant loss of ductility at the surface region of the high N-alloyed (1.0%) CrMnN steel 1.4452 at cryogenic temperatures. Furthermore, the strengthening of the considered alloys within the near-surface zone follows the expected trend, that hardening rate increases with decreasing SFE (1.4439 < 1.4301 < 1.4452). The concentration of the entire deformation energy within a 5 m layer leads to the high crack density within the worn surface of 1.4452.
Fig. 9. Effect of stress conditions on strain-induced phase transformation in the steel 1.4301.
3.2. Strain-induced-phase transformation All austenitic steels after sliding tests in LH2 were investigated concerning the strain-induced martensitic transformation by X-ray diffraction. Results displayed in Fig. 8 show that the austenitic microstructure in the CrNi steels 1.4439, 1.4876 and 1.4591 is not susceptible to phase transformation during wear processes. The stability of the austenite in these so-called stable alloys is not affected by increasing normal forces and decreasing sliding velocities. On the other hand, the metastable steel 1.4301 exhibits a significant martensitic transformation already after friction at RT. The influence of the loading and the sliding velocity on the phase transformation is shown in Fig. 9. As expected, an increased normal force leads to more intense martensitic transformation. In turn, slower sliding motions slightly increased the ␣-martensite fraction as well. This is in agreement with our previous investigations [20] by magnetic-inductive methods, where a significant variation of the martensite fraction with the sliding velocity could not be determined. Since a steady-state between the martensitic transformation and the wear rate is not achieved, as already reported during friction at RT [17], a lower martensite content at higher velocities can
Fig. 8. X-ray diffractrograms from the wear scars produced with: (a) FN = 5 N, v = 0.06 m/s and (b) FN = 10 N, v = 0.2 m/s.
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Fig. 10. X-ray diffractrogramm from the wear scar of the steel 1.4452.
be explained by a more severe removal of martensite from the contact zone at higher rotation frequencies. Contrary to the high-alloyed CrNi steels 1.4876 and 1.4591, where no martensitic transformation occurred at all, and the CrNi steel 1.4301 with an ␥/␣-transformation, in the CrMnN steel 1.4452 a ␥/-martensite transformation occurred after friction under FN = 10 N, v = 0.2 m/s (Fig. 10). The -martensite formed is strongly textured and very inhomogeneously distributed. The texture of the -martensite is due to the inheriting of the texture from the austenite, which was formed by plastic deformation in the wear scar.
Contrary to the steel 1.4452, in the high-alloyed steel 1.4876 with a high SFE, the hydrogen uptake markedly increases the SFE with increasing energies. Therefore, at low loadings, twinning plays an important role in the
3.3. Deformation mechanisms In order to clarify the mechanisms of cracking in LH2 and its variation with the loading conditions, the broadening of X-ray reflection profiles in the alloys 1.4452 and 1.4876 were quantitatively analyzed using a Rietveld refinement procedure regarding lattice constant (a), SFE, dislocation density (ρ) and twinning probability (β). Results are shown in Figs. 11 and 12. The different stress conditions were also displayed as the energy rate (Et ) introduced in the worn surface, which is given by Et = FN v. The hydrogen uptake can be visualized by means of an expansion of the unit cell. The lattice constant is determined by the Rietveld refinement procedure and represents a mean value for several reflections at ψ = 0◦ . In both alloys, a marked decrease of the lattice constant indicates that the hydrogen adsorption decreases with increasing loadings. In 1.4452, the SFE at 20 K is already very low, thus, no significant decrease of the SFE with hydrogen content can be observed. Increasing loadings, while keeping the sliding speed constant leads to higher plastic deformation and consequently to higher deformation rates on the microscale. The steep increase of twinning in the entire deformation process at high deformation energies accommodates deformation and internal stresses by higher strain rates at low temperature.
Fig. 11. Influence of the stress conditions on the hydrogen uptake, stacking fault energy and deformation mechanisms in the steel 1.4452.
Fig. 12. Influence of the stress conditions on the hydrogen uptake, stacking fault energy and deformation mechanisms in the steel 1.4876.
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Fig. 13. Heavily twinned zone within the nanostructured top surface layer of steel 1.4301 (a) and within the deformed near-surface zone (b) after friction in LH2 (FN = 5 N, v = 0.06 m/s).
deformation process and crack formation is observed. The low twin density at heavy stress conditions reflects the increased SFE and is responsible for the low crack density within the worn surface. Hydrogen uptake in austenitic steels takes place by hydrogen pair formation within the faulted zone [21–23]. Therefore, in alloys with high SFE, where the separation of the partial dislocations is very small, and thus, the faulted zones are very narrow, the hydrogen uptake may be hampered, especially if its release is favored by an intensified degradation of the surface zone at heavy loadings. In case of the alloys with low SFE, such as 1.4452, 1.4439 and 1.4301, hydrogen can be easily absorbed due to the generation of numerous lattice imperfections at the worn surface and the formation of extended faults. Thus, numerous intersecting slip bands, glide-twin interactions [17] (Fig. 13) as well as pile-ups at grain and phase boundaries within the nanostructured surface zone are able to initiate microcracks, especially if localized stress peaks by a hydrogen-induced pressure build-up. On the other hand, in alloys with a high SFE, a facilitated release of hydrogen takes place at high deformation energies. Hence, twinning, formation of slip bands and cracking appear only at low deformation energies.
4. Conclusions The chemical composition of austenitic steels determines the wear resistance at heavy loadings. Moderate N-additions enhance the stability of the austenitic microstructure and the wear performance at cryogenic temperatures. High N-contents, particularly in CrMn steels, induce a ductile to brittle transition in the top surface layer at extremely low temperatures, but the increased hardness determines a better wear resistance. Sliding friction in inert environments, such as air, LN2 and LHe, does not lead to crack formation in the contact zone. A hydrogen-induced embrittlement is not necessarily correlated to a reduced wear resistance. A low SFE matches the requirements for a good wear performance.
Extreme low SFE are connected to a loss of toughness at cryogenic temperatures. However, the wear resistance seems to be not affected. The performance during tribological stressing of austenitic alloys with high SFE can be clearly improved by small N-additions. High Ni-contents as well as its replacement by Mn in austenitic steels avoid the formation of the strain-induced ␣-martensite. In CrMn steels, a ␥/-phase transformation is possible with increasing loadings.
Acknowledgements The authors thank Professor Dr. W. Reimers, Technical University Berlin for the fruitful discussions. Dr. B. Hasse and Dr. T. Wroblewski are kindly acknowledged for technical support during the diffraction experiments at the beamline G3. This work was supported by the Deutsche Forschungsgemeinschaft (DFG) (Projects Py 9/4-1 and Hu 791/1-2).
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H. Pinto et al. / Wear 259 (2005) 424–431 [15] W. H¨ubner, H. Assmus, A. Pyzalla, H. Pinto, Fortschritte in der Metallographie, Sonderb¨ande der Praktischen Metallographie 35 (2004) 399–404. [16] H. Pinto, S. Stuke, A. Pyzalla, W. H¨ubner, K. Aßmus, TribologieFachtagung 2003, G¨ottingen, 22–24 September 2003, pp. 7/1– 7/10. [17] H. Pinto, A. Pyzalla, W. H¨ubner, K. Aßmus, DGM-Tagung Reibung und Verschleiß, F¨urth, 10–12 March 2004. [18] L.M. Robertson, Eng. Fract. Mech. 68 (2001) 671–692.
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[19] C. Zapffe, Am. Soc. Metals 39 (1947) 191–192. [20] K. Aßmus, W. H¨ubner, A. Pyzalla, H. Pinto, In: W.J. Bartz (Ed.), Proceedings of the 14th International Colloquium Tribology and Lubrication Engineering, vol. II, Technische Akademie Esslingen, 13–15 January 2004, pp. 1047–1054. [21] A.E. Pontini, J.D. Hermida, Scr. Mater. 37 (11) (1997) 1831–1837. [22] J.D. Hermida, A. Roviglione, Scr. Mater. 39 (8) (1998) 1145–1149. [23] A. Juan, L. Moro, G. Brizuela, E. Pronsato, Int. J. Hydrogen Energy 27 (2002) 333–338.
The effect of hydrogen on the deterioration of austenitic steels during wear at cryogenic temperature H. Pinto a,∗ , A. Pyzalla a , R. B¨uscher b , A. Fischer b , K. Aßmus c , W. H¨ubner c a
Vienna University of Technology, TU Wien, Institute for Materials Science and Technology (E308), Karlsplatz 13, A-1040 Vienna, Austria b University Duisburg-Essen, Ipe-Werkstofftechnik II, Lotharstreet 1, 47057 Duisburg, Germany c Federal Institute for Materials Research and Testing (BAM), Unter den Eichen 44-46, 12200 Berlin, Germany Received 16 August 2004; received in revised form 3 February 2005; accepted 4 February 2005 Available online 10 May 2005
Abstract Hydrogen represents an important alternative to fossil fuels. Hydrogen storage is possible as a gas, at room temperature (RT) at about 20 MPa pressure, and in a liquefied form, at cryogenic temperatures of about 20 K. The latter form is particularly attractive due to the possibility of stocking a large quantity of hydrogen within a small volume. In moving parts (e.g. of transport vehicles) cryogenic temperature and the presence of hydrogen strongly enhance wear processes and subsequently component failure. The present work deals with the deformation behaviour and the microstructural deterioration of austenitic CrNi- and CrMn high nitrogensteels during friction in liquid hydrogen at 20 K. The modified microstructure within the wear scar is studied by scanning electron microscopy and X-ray diffraction methods. Diffraction studies of wear scars reveal the importance of twinning during deformation at 20 K. This increase of twinning can be attributed to a hydrogen-induced reduction of stacking fault energy (SFE) in the austenitic steels. Interactions between twin boundaries and planar dislocation structures along with locally increased stresses led to the formation of extensive crack networks. The amount of hydrogen-induced surface cracks depends on the alloy composition and is not necessarily correlated to the wear resistance of the austenitic steels. © 2005 Elsevier B.V. All rights reserved. Keywords: Austenitic steel; Stacking fault energy; Hydrogen; Wear; Cryotechnology
1. Introduction Austenitic steels are materials commonly used in engines, which operate at cryogenic temperatures. These steels combine a high mechanical strength with sufficient ductility at extremely low temperatures. Cryogenic temperatures are encountered, e.g. in aerospace technology, where liquid hydrogen (LH2 ) became the standard propellant. Recently many efforts are being done to employ liquid hydrogen in fuel cells. Since hydrogen represents an important alternative to fossil fuels in the automotive sector, the production of hydrogen on a large scale along with its compatibility to the current technology of internal combustion engine are of great technological concern. Hydrogen storage is possible in tanks as a ∗
Corresponding author. Tel.: +43 1 58801 30836; fax: +43 1 58801 30897. E-mail address: [email protected] (H. Pinto).
0043-1648/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.wear.2005.02.057
gas, at room temperature (RT) at about 20 MPa pressure, and in a liquefied form, at cryogenic temperatures of about 20 K. The latter form is particularly attractive due to the possibility of stocking of a large quantity of hydrogen within a small volume. Due to the impracticable use of lubricants at extremely low temperatures, moving machine components are exposed to high wear rates. Wear processes are often correlated to an unexpected failure of sealings, valves and bearings. This could lead to an escape of hydrogen and the formation of an explosive atmosphere. Furthermore austenitic steels may undergo a deformationinduced martensitic-phase transformation [1], which is especially favoured by low temperatures. The formation of martensite within the worn surface enhances strength but also causes embrittlement and thus influences the degradation process of cryogenic machine components. Hence, the
H. Pinto et al. / Wear 259 (2005) 424–431
lifetime of tribo-systems strongly depends on the microstructural changes within the near-surface region. The aim of the present work is the study of the influence of the chemical composition and the tribological stress conditions on the stability of the austenitic microstructure, the deterioration process and the wear behaviour during friction in LH2 of five different austenitic steels, which are currently in use. They provide a continuous variation of the stacking fault energy (SFE) as well as of the austenite stability mainly due to their different Ni-, Mn- and N-contents. The investigations attempt to correlate the SFE (and its variation with the adsorbed hydrogen content at distinct loading conditions) to the microstructural modifications and the wear performance.
2. Experimental Five different austenitic steels were chosen for the present investigations. Their chemical compositions are shown in Table 1. The SFE at room temperature were determined as described in [2]. The austenitic steels chosen differ in their Cr, Ni, Mn and N contents. Thus, they have strongly different stacking fault energies. According to [3], the susceptibility to deformationinduced martensitic transformation increases with decreasing SFE. In [4,5], it has been shown that due to its influence on the active deformation mechanisms the SFE plays an important role in the wear behaviour. The other microstructural parameters, such as grain size and volume fraction of inclusions and ␦-ferrite, which also affect the mechanical strength of austenitic steels, did not vary considerably for the different alloys investigated. In order to simulate wear processes at cryogenic temperatures, ball-on-disc sliding friction tests have been carried out using a special cryotribometer [6] in a liquid hydrogen LH2 (20 K) atmosphere at the Federal Institute for Materials Research and Testing (BAM) in Berlin, Germany. The loading parameters for the sliding tests were: normal force FN = 5 and 10 N, sliding speed, v = 0.06 and 0.2 m/s and sliding distance s = 1000 m. An aluminium oxide ceramic ball was employed as counterbody in order to minimize the transfer layer between the two sliding bodies. The worn surfaces were inspected using optical and scanning electron microscopy (SEM). The wear resistance was evaluated by means of the worn area on the ball specimen using SEM. Analysis of loss of weight and scar depth led to
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very high uncertainties due to the small scar dimensions. In order to determine the depth of the deformed near-surface layer, microhardness tests HV0.01 at the CrNi steel samples and nanohardness tests using a Berkovich identer at the CrMn steel samples were carried out. Nanohardness measurements were necessary due to extremely thin layers, which were formed on the surface in the steel 1.4452. For the nanohardness tests, a penetration depth of 0.3 m was chosen. Synchrotron X-ray diffraction studies were performed at the experimental station G3, at HASYLAB, DESY, in Hamburg, Germany, using a wavelength similar to Co K␣radiation. The strain-induced martensite formation within the wear scar, the deformation-induced lattice defects were quantified. For the evaluation of the martensite content after friction, the texture in each phase was taken into account [7,8]. The recorded diffractograms from the wear scar have been subjected to a Rietveld refinement procedure using the freeware material analysis using diffraction (MAUD) [9–11] in order to analyze the deformed microstructure. The instrument-broadening function has been determined by means of a standard silicon powder with a well-known grain size distribution. The SFE and the dislocation densities were evaluated from the microstrains and stacking fault probabilities determined by the Rietveld refinement procedure [4,12], as follows: K111 ω0 G(1 1 1) a0 A−0.37 ε250 1 1 1 SFE = (1) √ α π 3 where ε250 1 1 1 is the microstrain along the 1 1 1 crystallographic direction, α the stacking fault probability, K111 ω0 the proportionality constant, G(1 1 1) the shear modulus in the (1 1 1) fault plane = 1/3(c44 + c11 − c12 ), a0 the lattice constant, A the Zener anisotropy = 2c44 /(c11 − c12 ) and cij is the elastic stiffness coefficient. ρ=
48πα(SFE) (Ke + Ks )a02 d1 1 1
(2)
where ρ is the dislocation density, and d1 1 1 is the interplanar spacing and c 1/2 44 (c11 − c12 ) Ks = 2 and 1/2 c44 (c11 − c12 ) . Ke = (c11 + c12 ) c11 (c11 + c12 + 2c44 )
Table 1 Chemical composition of the investigated materials Materiala
SFERT (mJ/m2 )
Cr
Ni
Mn
Mo
C
N
Fe
1.4452 1.4301 1.4439 1.4876 1.4591
12 21 30 42 71
16.0–20.0 17.0–19.5 16.5–18.5 20.5 32.85
– 8.0–10.5 12.5–14.5 30.65 30.95
12.0–16.0 2.0 1.5 0.69 –
2.5–4.2 – 4.0–4.5 – 1.67
<0.15 <0.07 0.03 0.061 0.007
0.75–1.0 <0.11 0.12–0.22 – 0.39
Bal. Bal. Bal. Bal. Bal.
a
German steel grades.
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Fig. 1. Effect of increasing sliding velocities and normal forces on the wear resistance.
3. Results and discussion 3.1. Wear behaviour The wear performance during sliding friction tests in LH2 was evaluated by means of the wear scars created on the counterbody. These worn areas produced on the Al2 O3 -balls were chosen for the evaluation of the wear amount, since on the steel specimens, the scar depths are very small (comparable to the roughness of the polished surface) and the scar widths may lead to an erroneous interpretation due to the material transfer between disc and counterbody during adhesive wear. Fig. 1 displays the variation of the scar widths on the ball specimen with the tribological stress conditions for the different austenitic steels. The error bars give the standard deviation between the diameters of the scars measured along and perpendicular to the sliding direction. Results reveal that at mild tribological stressing (FN = 5 N, v = 0.06 m/s) the wear resistance of all investigated alloys is nearly the same. An increase of the sliding speed (FN = 5 N, v = 0.2 m/s) leads to a reduction of wear in all N-alloyed steels and to a linear increase of wear in the N-free 1.4876. On
the other hand, distinct wear performances for each investigated austenitic steel appear at more severe stress conditions (FN = 10 N, v = 0.2 m/s). The metastable CrNi steel 1.4301 and the CrMnN-steel 1.4452 with the lowest SFE present the best wear resistance The second best wear performance is found for the Mo-alloyed steel 1.4439 with about 0.17% N and moderate SFE (Fig. 1). For these three steel grades (1.4301, 1.4452 and 1.4439) the increase of normal force and sliding speed does not affect the material removal. A high SFE leads clearly to a loss of wear resistance. The highalloyed steel 1.4591 with 0.39% N exhibits an intermediate wear performance, whilst the Ni-free steel 1.4876 appears with the broadest scar. These distinct wear behaviours suggest an important role of the strain hardening by martensite formation on the wear performance of the metastable alloy 1.4301. In addition a positive effect of Mo- and N-alloying in CrNi steels can be observed. Moderate Mo- and N-contents are known to increase the yield strength, while sufficient toughness remains. This explains the enhanced wear resistance at cryogenic temperatures. A high N-content (1.0%) in the CrMn steel can result in loss of ductility [13]. However, the increased mechanical strength of the worn surface layer may be responsible for the best wear resistance of the CrMnN steel 1.4452. Fig. 2 shows typical images of the topography of the wear scars in inert environments. By an adhesive wear mechanism, particles are detached from the tribologically modified zone, then transferred to the counterbody and finally, re-transferred to the sample surface. The wear particles are generated as soon as the mechanical strength reaches its limit during the shear deformation of the surface. Crack formation could not be detected even at the lowest temperature, i.e. LHe (4.2 K), in all investigated alloys. In liquid hydrogen, the deformation behaviour in the contact zone seems to be different, as shown in Figs. 3–5. Crack formation is observed in particles deposited as well as in the worn matrix. The cracking density depends on the alloy and the loading conditions. In our previous works [14–17], we reported that an appreciable decrease of the SFE along with a very high yield strength and texture hardening favour cracking at the worn
Fig. 2. Wear scars after friction in LHe (FN = 5 N, v = 0.2 m/s): (a) steel 1.4301 and (b) steel 1.4439.
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Fig. 3. Steel 1.4452 after testing in LH2 at different loading conditions.
Fig. 4. Steel 1.4876 after testing in LH2 at different loading conditions.
Fig. 5. Steel 1.4591 after testing in LH2 at different loading conditions.
surface during sliding in LH2 by the pressure build-up mechanism [18,19]. In the present work, systematic SEM studies of the wear scars were carried out in order to evaluate the effect of the stress conditions, and the chemical composition on the cracking process for the different austenitic alloys. Figs. 3–5 display the worn surface of the austenitic steels 1.4452, 1.4876 and 1.4591 after friction at different conditions. In 1.4452, cracking occurs at all loading conditions and is intensified by increasing loadings. In turn, only very few cracks are formed in 1.4876 at the slightest applied stress condition of 5 N and 0.06 m/s. Heavier loading does not increase the crack formation but rather suppresses it. The steel 1.4591 at 5 N and 0.06 m/s is more sensitive to cracking than 1.4876 at 5 N and 0.06 m/s but similar to the later presents reduced cracking at heavier stress conditions. This suggests that in case of 1.4876 and 1.4591, the more pronounced deformation and severe loading conditions do not enhance the hydrogen uptake but likely facilitates the hydrogen release. It is assumed that hydrogen is continuously driven out and is not able to be soluted in the
austenitic crystal lattice. In all samples, the observed cracks formed within a nanocrystalline debris layer on the worn surface (Fig. 6). Therefore, crack appearance is not necessarily connected to a low wear resistance at cryogenic temperatures. As shown
Fig. 6. Deteriorated contact layer and crack network in the steel 1.4301 after friction in LH2 (FN = 5 N, v = 0.06 m/s).
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Fig. 7. Hardness–depth profiles of the near-surface zones after sliding in LH2 with FN = 5 N and v = 0.06 m/s.
in Fig. 1, 1.4876 and 1.4591 exhibit unsatisfactory wear performances but are less embrittled in presence of hydrogen. On the contrary, the metastable 1.4301, the steel 1.4439 and even the steel 1.4452 with the best wear performance reveal a very high crack density. The in-depth extension of the strengthened near-surface zone in the different austenitic alloys after friction in LH2 was evaluated by means of micro- and nanohardness tests. Fig. 7 shows the hardness–depth profiles for the steels 1.4452, 1.4301 and 1.4439 after tribological stressing with 5 N and 0.06 m/s. The depth of the hardened zone amounts to about 50 m for both CrNi steels (1.4301 and 1.4439), but the hardened zone is clearly thinner for the CrMnN steel 1.4452 (approximately 5 m). This reflects a pronounced increase of mechanical strength accompanied by a significant loss of ductility at the surface region of the high N-alloyed (1.0%) CrMnN steel 1.4452 at cryogenic temperatures. Furthermore, the strengthening of the considered alloys within the near-surface zone follows the expected trend, that hardening rate increases with decreasing SFE (1.4439 < 1.4301 < 1.4452). The concentration of the entire deformation energy within a 5 m layer leads to the high crack density within the worn surface of 1.4452.
Fig. 9. Effect of stress conditions on strain-induced phase transformation in the steel 1.4301.
3.2. Strain-induced-phase transformation All austenitic steels after sliding tests in LH2 were investigated concerning the strain-induced martensitic transformation by X-ray diffraction. Results displayed in Fig. 8 show that the austenitic microstructure in the CrNi steels 1.4439, 1.4876 and 1.4591 is not susceptible to phase transformation during wear processes. The stability of the austenite in these so-called stable alloys is not affected by increasing normal forces and decreasing sliding velocities. On the other hand, the metastable steel 1.4301 exhibits a significant martensitic transformation already after friction at RT. The influence of the loading and the sliding velocity on the phase transformation is shown in Fig. 9. As expected, an increased normal force leads to more intense martensitic transformation. In turn, slower sliding motions slightly increased the ␣-martensite fraction as well. This is in agreement with our previous investigations [20] by magnetic-inductive methods, where a significant variation of the martensite fraction with the sliding velocity could not be determined. Since a steady-state between the martensitic transformation and the wear rate is not achieved, as already reported during friction at RT [17], a lower martensite content at higher velocities can
Fig. 8. X-ray diffractrograms from the wear scars produced with: (a) FN = 5 N, v = 0.06 m/s and (b) FN = 10 N, v = 0.2 m/s.
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Fig. 10. X-ray diffractrogramm from the wear scar of the steel 1.4452.
be explained by a more severe removal of martensite from the contact zone at higher rotation frequencies. Contrary to the high-alloyed CrNi steels 1.4876 and 1.4591, where no martensitic transformation occurred at all, and the CrNi steel 1.4301 with an ␥/␣-transformation, in the CrMnN steel 1.4452 a ␥/-martensite transformation occurred after friction under FN = 10 N, v = 0.2 m/s (Fig. 10). The -martensite formed is strongly textured and very inhomogeneously distributed. The texture of the -martensite is due to the inheriting of the texture from the austenite, which was formed by plastic deformation in the wear scar.
Contrary to the steel 1.4452, in the high-alloyed steel 1.4876 with a high SFE, the hydrogen uptake markedly increases the SFE with increasing energies. Therefore, at low loadings, twinning plays an important role in the
3.3. Deformation mechanisms In order to clarify the mechanisms of cracking in LH2 and its variation with the loading conditions, the broadening of X-ray reflection profiles in the alloys 1.4452 and 1.4876 were quantitatively analyzed using a Rietveld refinement procedure regarding lattice constant (a), SFE, dislocation density (ρ) and twinning probability (β). Results are shown in Figs. 11 and 12. The different stress conditions were also displayed as the energy rate (Et ) introduced in the worn surface, which is given by Et = FN v. The hydrogen uptake can be visualized by means of an expansion of the unit cell. The lattice constant is determined by the Rietveld refinement procedure and represents a mean value for several reflections at ψ = 0◦ . In both alloys, a marked decrease of the lattice constant indicates that the hydrogen adsorption decreases with increasing loadings. In 1.4452, the SFE at 20 K is already very low, thus, no significant decrease of the SFE with hydrogen content can be observed. Increasing loadings, while keeping the sliding speed constant leads to higher plastic deformation and consequently to higher deformation rates on the microscale. The steep increase of twinning in the entire deformation process at high deformation energies accommodates deformation and internal stresses by higher strain rates at low temperature.
Fig. 11. Influence of the stress conditions on the hydrogen uptake, stacking fault energy and deformation mechanisms in the steel 1.4452.
Fig. 12. Influence of the stress conditions on the hydrogen uptake, stacking fault energy and deformation mechanisms in the steel 1.4876.
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Fig. 13. Heavily twinned zone within the nanostructured top surface layer of steel 1.4301 (a) and within the deformed near-surface zone (b) after friction in LH2 (FN = 5 N, v = 0.06 m/s).
deformation process and crack formation is observed. The low twin density at heavy stress conditions reflects the increased SFE and is responsible for the low crack density within the worn surface. Hydrogen uptake in austenitic steels takes place by hydrogen pair formation within the faulted zone [21–23]. Therefore, in alloys with high SFE, where the separation of the partial dislocations is very small, and thus, the faulted zones are very narrow, the hydrogen uptake may be hampered, especially if its release is favored by an intensified degradation of the surface zone at heavy loadings. In case of the alloys with low SFE, such as 1.4452, 1.4439 and 1.4301, hydrogen can be easily absorbed due to the generation of numerous lattice imperfections at the worn surface and the formation of extended faults. Thus, numerous intersecting slip bands, glide-twin interactions [17] (Fig. 13) as well as pile-ups at grain and phase boundaries within the nanostructured surface zone are able to initiate microcracks, especially if localized stress peaks by a hydrogen-induced pressure build-up. On the other hand, in alloys with a high SFE, a facilitated release of hydrogen takes place at high deformation energies. Hence, twinning, formation of slip bands and cracking appear only at low deformation energies.
4. Conclusions The chemical composition of austenitic steels determines the wear resistance at heavy loadings. Moderate N-additions enhance the stability of the austenitic microstructure and the wear performance at cryogenic temperatures. High N-contents, particularly in CrMn steels, induce a ductile to brittle transition in the top surface layer at extremely low temperatures, but the increased hardness determines a better wear resistance. Sliding friction in inert environments, such as air, LN2 and LHe, does not lead to crack formation in the contact zone. A hydrogen-induced embrittlement is not necessarily correlated to a reduced wear resistance. A low SFE matches the requirements for a good wear performance.
Extreme low SFE are connected to a loss of toughness at cryogenic temperatures. However, the wear resistance seems to be not affected. The performance during tribological stressing of austenitic alloys with high SFE can be clearly improved by small N-additions. High Ni-contents as well as its replacement by Mn in austenitic steels avoid the formation of the strain-induced ␣-martensite. In CrMn steels, a ␥/-phase transformation is possible with increasing loadings.
Acknowledgements The authors thank Professor Dr. W. Reimers, Technical University Berlin for the fruitful discussions. Dr. B. Hasse and Dr. T. Wroblewski are kindly acknowledged for technical support during the diffraction experiments at the beamline G3. This work was supported by the Deutsche Forschungsgemeinschaft (DFG) (Projects Py 9/4-1 and Hu 791/1-2).
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