- Email: [email protected]
Interplay of cold working and nitrogen diffusion in austenitic stainless steel
NIM B Beam Interactions with Materials & Atoms
Nuclear Instruments and Methods in Physics Research B 257 (2007) 442–446 www.elsevier.com/locate/nimb
Interplay of cold working and nitrogen diffusion in austenitic stainless steel D. Manova
b
a,*
, I.-M. Eichentopf a, S. Heinrich a, S. Ma¨ndl a, E. Richter b, H. Neumann a, B. Rauschenbach a
a Leibniz-Institut fu¨r Oberfla¨chenmodifizierung, Permoserstr. 15, 04318 Leipzig, Germany Forschungszentrum Rossendorf, Institut fu¨r Ionenstrahlphysik und Materialforschung, 01314 Dresden, Germany
Available online 11 February 2007
Abstract For hardness measurements, the indentation depth is about 10% of the effective information depth under static loading. A change of the wear mechanism is observed under lateral loading conditions in oscillating ball-on-disc tests for nitrided austenitic stainless steel with an expanded austenite surface layer from abrasive to a subsurface plastic flow with a redistribution of the inserted nitrogen. This leads to an effective nitriding depth about 3–5 larger than the actual nitrided zone. 2007 Elsevier B.V. All rights reserved. PACS: 52.77.Dq; 81.70.Bt; 81.65.Lp; 66.30. h Keywords: PIII; Ion implantation; Steel; Diffusion; Wear
1. Introduction Nitriding of austenitic stainless steel at temperatures between 300 and 400 C results in formation of a modified layer with outstanding hardness up to 1200 HV together with high wear resistance, while the excellent corrosion resistance is still preserved. This layer is characterised by nitrogen concentration up to 30% on the surface and a lattice expansion of 5–10% [1–3]. It is well known that nitrogen diffusion in austenitic stainless steel is a complicated process, not fully understood and still open to debate. However the microstructure is assumed to play a significant role, as martensitic and ferritic stainless steels do not show the distinct acceleration of diffusion [4] and nanocrystalline austenitic stainless steel shows a diffusion rate about 1000 times faster than the microcrystalline modification [5]. The higher hardness and lower wear make these modified surfaces ideal candidates for wear protection under
*
Corresponding author. Tel.: +49 341 235 3606; fax: + 49 341 235 2313. E-mail address: [email protected] (D. Manova).
0168-583X/$ - see front matter 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.nimb.2007.01.257
normal loading. The higher hardness i.e. larger elastic modulus is effective for indentation depths about 10% of layer thickness [6]. Nearly all investigations of wear behaviour of nitrided layers show very low wear rates for prolonged times without any significant increase at the end of the experiment [7–9]. As the wear rate decreases by 2–3 orders of magnitude, it may be assumed that all these experiments stop the wear while still within the layer. Increasing of wear resistance by factor of 200 for the surface of austenitic stainless steel after nitriding was shown by this group [10]. A reduction of this factor from 200 to 75 is observed near the end of the experiment well beyond of the implanted layer thickness. Open questions are about a dynamic hardening mechanism during wear test and any change of the wear mode due to nitriding. In this presentation, detailed investigations of austenitic stainless steel samples nitrided by plasma immersion ion implantation (PIII) and low energy implantation (LEI) are performed during and after wear test. A comparison with base material was made. A correlation of the results from metallographic cross-sections with hardness and wear test parameters allows an inside into the hardening
D. Manova et al. / Nucl. Instr. and Meth. in Phys. Res. B 257 (2007) 442–446
mechanism and tailoring the modified properties deep into the bulk material.
443
substrate at a load of 100 and 200 mN and additional lateral forces. The resulting average speed was 0.02 mm/s at 1.5 GPa contact pressure as total wear path of 4000 lm was reached at track length of 50 lm. Subsequently metallographic cross-sections were prepared from nitrided and non-implanted samples after wear test with WC ball and investigated with scanning electron microscopy (SEM). The etching solution was aqua regia. Some additional images were obtained from optical microscope.
2. Experiment Nitrogen was inserted by PIII and LEI into polished to mirror finished coupons from austenitic stainless steel grade X5CrNi18.10 (AISI 304/DIN 1.4301). The implantations were performed in a high vacuum chamber with a base pressure less 5 · 10 7 mbar equipped with an ECR plasma source. Ten kilovolt of negative voltage pulses, pulse length 15 ls and rise time less than 0.5 ls were applied to the samples immersed in a nitrogen plasma. Different pulse frequencies were used to maintain the process temperature between 300 and 400 C, with the temperature measured using a pyrometer calibrated against a thermocouple [11]. No additional heating or cooling was employed. The total implanted fluence was up to 1018 at/ cm2 for implantation time of up to 60 min. The same high vacuum chamber was used for LEI nitriding with a broad beam ECR ion at 1.5 kV. The samples were heated up to the process temperature between 300 and 400 C and nitriding only began at the desired temperature. The implanted fluence was between 1 and 2 · 1018 at/cm2. The thicknesses of the nitrided layers were obtained from SIMS measurements. Dynamic hardness measurements were carried out using Berkovich indenter at a load of 20 and 50 mN. No corrections for the elastic recovery was made therefore the real values should be between 15% and 21% higher. Wear test measurements were performed using an oscillating dry ball-on-disc geometry against 3 mm WC ball. Three Newton load was applied resulting in a Hertzian contact pressure of 1 GPa. The average speed was 15 mm/s for a track length 2 mm and wear paths up to 400 m. Additional wear data were collected from nanohardness equipment using a 10 lm diamond ball pressed on the
3. Results Fig. 1 shows hardness data for nitrided stainless steel samples together with non-nitrided as a function of layer thickness. Implantations were made using PIII and LEI at temperatures in the range of 300–400 C. As different process temperatures were chosen nitrided layers with thickness in the range from 500 nm to 8 lm were produced. Two different loads of 20 and 50 mN were used for untreated and LEI nitrided samples, while for PIII implanted samples only data for lower load of 20 mN are shown. The hardness of the initial surface layer can be inferred to 1100–1200 HV, while the measured values are a mixture of the surface layer and the bulk. The value itself depends on their relative contribution width on information depth about 10 times larger than the indentation depth. Wear tests were carried out for the above investigated sample using the equipment for the nanohardness measurements and a diamond ball with 10 lm diameter applying two different normal loads of 100 and 200 mN and an additional lateral force to keep the speed constant. The results extracted from the wear data are shown in Fig. 2, where the wear rates deduced from the measurements are presented as a function of the layer thickness. Up to 2 orders of
15 14 13 12
Hardness (GPa)
11 10 9 8
Load untreated LEI PIII
20 mN 50 mN
Temperature 300 - 400 ˚C LEI 15 min, PIII 60 min
7 6 5 4 3 2 1 0
0.00
1
1
0
Layer Thickness (µm) Fig. 1. Effective surface hardness as a function of layer thickness. Please note the logarithmic depth scale.
444
D. Manova et al. / Nucl. Instr. and Meth. in Phys. Res. B 257 (2007) 442–446
Fig. 2. Wear rate as a function of layer thickness. Please note the logarithmic depth scale.
magnitude lower wear rates are observed after nitriding. However, in contrast to the thickness dependent hardness, nearly no influence of the layer thickness on the wear rates at low loads is seen for both methods of implantations. At high loads and thin layers, a different breakthrough is observed, which is not subject of this paper. The low wear, independent of layer thickness is now investigated further. Fig. 3 depicts one-dimensional profiles across wear tracks after different total length on two samples, non-implanted and implanted. The thickness of the nitrided layer was about 1.5 lm, as obtained from secondary ion mass spectroscopy measurements. Using a larger ball (B 3 mm) at about the same contact pressure leads to quantifiable wear tracks in the wear tests. A strong decrease of the wear track depth for the same total wear
path compared to the control sample is observed for the nitrided sample. Additional analysing of the 3D profiles yields a wear rate decrease by about a factor of 200 after implantation. However only slight increase in the wear rate of the nitrided sample of about factor of 3 is detected after total wear path length corresponding to 5–10 lm, which is far beyond the nitrided layer. The profiles for both samples are geometrically similar with the width given by contact area, which increases with the total wear path. This can be seen especially very good for the implanted sample after shorter wear test. Redeposition of material during the wear test is clearly seen for the non-treated sample, while it is nearly missing for the nitrided sample, not only at the edges but also at the reversal points.
Fig. 3. One-dimensional wear track profiles after different wear paths (different time for wear test) for untreated and nitrogen implanted (10 kV, 350 C, 2 h) samples.
D. Manova et al. / Nucl. Instr. and Meth. in Phys. Res. B 257 (2007) 442–446
Fig. 4. SEM micrographs in the vicinity of a wear track from cross section of untreated (a) and nitrided sample (b).
Metallographic cross-sections were made before and after wear test on implanted and the as-received samples. Wear tracks for this investigation were manufactured with the WC ball due to the reason mentioned above. Microscopic viewgraphs of the area in the vicinity of the wear tracks, obtained from SEM, are presented in Fig. 4(a) non-implanted sample and (b) a sample after nitrogen implantation. The microstructure just below the wear track on the non-treated sample shows strongly deformed grains in the direction perpendicular to the applied load and indicates abrasive mode after the shear strength is reached. In contrast, the implanted sample shows a microstructure with a drastic change of grain shape in the vicinity of the wear track, as the individual grains are barely visible and the structure resembles highly deformed material with strongly elongated grains. Plan view images of the samples, after wear experiments performed with a 10 lm diamond ball, are shown in Fig. 5, as obtained from optical microscope. Formation of a large pile-up and a huge number of dislocations is clearly visible in Fig. 5(a) presenting the microstructure of the nonimplanted sample. In contrast, the observable pile-up is very small and no visible dislocations can be found near the wear track for the implanted material (Fig. 5(b)).
445
Fig. 5. Plan view optical micrographs of a wear track on untreated (a) and nitrided (b) sample.
4. Discussion The presented results can be used to highlight the different deformation mechanism for the base material and for nitrided surfaces. For the non-implanted sample the applied load results in an agglomeration of dislocations leading to grain refinement with much smaller gain size in the area where the local stress exceeds the strength of stainless steel. The local pressure just below the contact area exceeding 1.0 GPa has to be put in relation to the strength of about 0.5 GPa [12]. A significant pile-up after wear tests on untreated austenitic stainless steel, formed by particles removed from the wear track, indicates abrasive wear correlated with ploughing into the surface by the counterbody ball. In contrast, the same applied load on the nitrided samples leads only to strong deformation of grains resulting in highly anisotropic thin fibres with less pronounced dislocations. The nearly missing pile-up for the nitrided sample, together with a non-existent weight gain of the WC balls, indicates a completely different wear mechanism. As the penetration depth of the ball at the same contact pressure is smaller for a higher hardness, respective elastic modulus of the surface, a transition from surface removal
446
D. Manova et al. / Nucl. Instr. and Meth. in Phys. Res. B 257 (2007) 442–446
to deformation due to a indentation below the threshold where cutting occurs can be envisaged [13]. For shallower cutting depths and the same tool geometry, only substantial subsurface plastic flow will occur instead of chip formation. However, an additional effect is proposed to transform the surface properties to explain the complete observations, including the missing dislocations within the nitrided surfaces. Cracking and abrasion are indicative of large intergranular stresses, in contrast to elastic intragranular stored energy [14], which can be correlated with the stacking fault energies [15]. Similar to high nitrogen steels containing molybdenum, a decreased stacking fault energy there results in a decreased fatigue crack growth rate there [16]. As only 0.2 wt% nitrogen were necessary to observe this effect, an even more pronounced manifestation can be expected at around 15–25 at% in the present case. As a consequence, the nitrided surface layer is not removed by the wear process but recycled and folded into the subsurface below the wear track during the actual wear process – similar to cold work hardening but more effective – where an additional wear resistant layer is formed dynamically during the wear process. This leads to an effective wear resistant layer about 3–5 times thicker than implanted zone.
5. Summary and conclusions Very hard and wear resistant layers are formed after nitrogen insertion into austenitic stainless steel using PIII and LEI. An outstanding wear resistance was observed for an effective layer thickness of about 3–5 times the thickness of the nitrided layer. A change from abrasive wear towards substantial subsurface plastic flow is mediated by the higher elastic modulus and the modified stacking fault energy of the expanded austenite. For applications this means a wear protection of the nitrided surface extends beyond the original layer thickness without thermal or
stress induced experiments.
nitrogen
diffusion
during
the
wear
Acknowledgements Parts of this work were supported by European Regional Development Fund (ERDF) and the State of Saxony (Project 106967/1650), as well as the German Federation of Industrial Research Associations (AiF, Project KF 0189601UK5). References [1] C. Blawert, B.L. Mordike, Y. Jira´skova´, O. Schneeweiss, Surf. Coat. Technol. 116–119 (1999) 189. [2] D.L. Williamson, J.A. Davis, P.J. Wilbur, J.J. Vajo, R. Wei, J.N. Matossian, Nucl. Instr. and Meth. B 127–128 (1997) 930. [3] G.A. Collins, R. Hutchings, K.T. Short, J. Tendys, Surf. Coat. Technol. 103–104 (1998) 212. [4] D. Manova, I.-M. Eichentopf, D. Hirsch, S. Ma¨ndl, H. Neumann, B. Rauschenbach, IEEE Trans. Plasma Sci. 34 (2006) 1136. [5] W.P. Tong, N.R. Tao, Z.B. Wang, J. Lu, K. Lu, Science 299 (2003) 686. [6] H. Bu¨ckle, Mikroha¨rtepru¨fung und ihre Anwendung, Berliner Union Verlag, Stuttgart, 1957. [7] M. Samandi, B.A. Shedden, D.I. Smith, G.A. Collins, R. Hutchings, J. Tendys, Surf. Coat. Technol. 59 (1993) 261. [8] E. Menthe, K.-T. Rie, Surf. Coat. Technol. 116–119 (1999) 199. [9] S. Ma¨ndl, R. Gu¨nzel, E. Richter, W. Mo¨ller, Surf. Coat. Technol. 100–101 (1998) 372. [10] D. Manova, S. Ma¨ndl, H. Neumann, B. Rauschenbach, Surf. Coat. Technol. 200 (2005) 137. [11] D. Manova, S. Ma¨ndl, B. Rauschenbach, Plasma Sources Sci. Technol. 10 (2001) 423. [12] L. Vitos, P.A. Korzhavyi, B. Johansson, Adv. Eng. Mater. 4 (2004) 228. [13] J.D. Thiele, S.N. Melkote, J. Manuf. Proc. 2 (2000) 270. [14] Y.-D. Wang, H. Tian, A.D. Stoica, X.-L. Wang, P.K. Liaw, J.W. Richardson, Nat. Mater. 2 (2003) 101. [15] A.G. Froseth, P.M. Derlet, H. Van Swygenhoven, Adv. Eng. Mater. 7 (2005) 16. [16] M. Murayama, K. Hono, H. Hirukawa, T. Ohmura, S. Matsuoka, Scripta Mater. 41 (1999) 467.
Nuclear Instruments and Methods in Physics Research B 257 (2007) 442–446 www.elsevier.com/locate/nimb
Interplay of cold working and nitrogen diffusion in austenitic stainless steel D. Manova
b
a,*
, I.-M. Eichentopf a, S. Heinrich a, S. Ma¨ndl a, E. Richter b, H. Neumann a, B. Rauschenbach a
a Leibniz-Institut fu¨r Oberfla¨chenmodifizierung, Permoserstr. 15, 04318 Leipzig, Germany Forschungszentrum Rossendorf, Institut fu¨r Ionenstrahlphysik und Materialforschung, 01314 Dresden, Germany
Available online 11 February 2007
Abstract For hardness measurements, the indentation depth is about 10% of the effective information depth under static loading. A change of the wear mechanism is observed under lateral loading conditions in oscillating ball-on-disc tests for nitrided austenitic stainless steel with an expanded austenite surface layer from abrasive to a subsurface plastic flow with a redistribution of the inserted nitrogen. This leads to an effective nitriding depth about 3–5 larger than the actual nitrided zone. 2007 Elsevier B.V. All rights reserved. PACS: 52.77.Dq; 81.70.Bt; 81.65.Lp; 66.30. h Keywords: PIII; Ion implantation; Steel; Diffusion; Wear
1. Introduction Nitriding of austenitic stainless steel at temperatures between 300 and 400 C results in formation of a modified layer with outstanding hardness up to 1200 HV together with high wear resistance, while the excellent corrosion resistance is still preserved. This layer is characterised by nitrogen concentration up to 30% on the surface and a lattice expansion of 5–10% [1–3]. It is well known that nitrogen diffusion in austenitic stainless steel is a complicated process, not fully understood and still open to debate. However the microstructure is assumed to play a significant role, as martensitic and ferritic stainless steels do not show the distinct acceleration of diffusion [4] and nanocrystalline austenitic stainless steel shows a diffusion rate about 1000 times faster than the microcrystalline modification [5]. The higher hardness and lower wear make these modified surfaces ideal candidates for wear protection under
*
Corresponding author. Tel.: +49 341 235 3606; fax: + 49 341 235 2313. E-mail address: [email protected] (D. Manova).
0168-583X/$ - see front matter 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.nimb.2007.01.257
normal loading. The higher hardness i.e. larger elastic modulus is effective for indentation depths about 10% of layer thickness [6]. Nearly all investigations of wear behaviour of nitrided layers show very low wear rates for prolonged times without any significant increase at the end of the experiment [7–9]. As the wear rate decreases by 2–3 orders of magnitude, it may be assumed that all these experiments stop the wear while still within the layer. Increasing of wear resistance by factor of 200 for the surface of austenitic stainless steel after nitriding was shown by this group [10]. A reduction of this factor from 200 to 75 is observed near the end of the experiment well beyond of the implanted layer thickness. Open questions are about a dynamic hardening mechanism during wear test and any change of the wear mode due to nitriding. In this presentation, detailed investigations of austenitic stainless steel samples nitrided by plasma immersion ion implantation (PIII) and low energy implantation (LEI) are performed during and after wear test. A comparison with base material was made. A correlation of the results from metallographic cross-sections with hardness and wear test parameters allows an inside into the hardening
D. Manova et al. / Nucl. Instr. and Meth. in Phys. Res. B 257 (2007) 442–446
mechanism and tailoring the modified properties deep into the bulk material.
443
substrate at a load of 100 and 200 mN and additional lateral forces. The resulting average speed was 0.02 mm/s at 1.5 GPa contact pressure as total wear path of 4000 lm was reached at track length of 50 lm. Subsequently metallographic cross-sections were prepared from nitrided and non-implanted samples after wear test with WC ball and investigated with scanning electron microscopy (SEM). The etching solution was aqua regia. Some additional images were obtained from optical microscope.
2. Experiment Nitrogen was inserted by PIII and LEI into polished to mirror finished coupons from austenitic stainless steel grade X5CrNi18.10 (AISI 304/DIN 1.4301). The implantations were performed in a high vacuum chamber with a base pressure less 5 · 10 7 mbar equipped with an ECR plasma source. Ten kilovolt of negative voltage pulses, pulse length 15 ls and rise time less than 0.5 ls were applied to the samples immersed in a nitrogen plasma. Different pulse frequencies were used to maintain the process temperature between 300 and 400 C, with the temperature measured using a pyrometer calibrated against a thermocouple [11]. No additional heating or cooling was employed. The total implanted fluence was up to 1018 at/ cm2 for implantation time of up to 60 min. The same high vacuum chamber was used for LEI nitriding with a broad beam ECR ion at 1.5 kV. The samples were heated up to the process temperature between 300 and 400 C and nitriding only began at the desired temperature. The implanted fluence was between 1 and 2 · 1018 at/cm2. The thicknesses of the nitrided layers were obtained from SIMS measurements. Dynamic hardness measurements were carried out using Berkovich indenter at a load of 20 and 50 mN. No corrections for the elastic recovery was made therefore the real values should be between 15% and 21% higher. Wear test measurements were performed using an oscillating dry ball-on-disc geometry against 3 mm WC ball. Three Newton load was applied resulting in a Hertzian contact pressure of 1 GPa. The average speed was 15 mm/s for a track length 2 mm and wear paths up to 400 m. Additional wear data were collected from nanohardness equipment using a 10 lm diamond ball pressed on the
3. Results Fig. 1 shows hardness data for nitrided stainless steel samples together with non-nitrided as a function of layer thickness. Implantations were made using PIII and LEI at temperatures in the range of 300–400 C. As different process temperatures were chosen nitrided layers with thickness in the range from 500 nm to 8 lm were produced. Two different loads of 20 and 50 mN were used for untreated and LEI nitrided samples, while for PIII implanted samples only data for lower load of 20 mN are shown. The hardness of the initial surface layer can be inferred to 1100–1200 HV, while the measured values are a mixture of the surface layer and the bulk. The value itself depends on their relative contribution width on information depth about 10 times larger than the indentation depth. Wear tests were carried out for the above investigated sample using the equipment for the nanohardness measurements and a diamond ball with 10 lm diameter applying two different normal loads of 100 and 200 mN and an additional lateral force to keep the speed constant. The results extracted from the wear data are shown in Fig. 2, where the wear rates deduced from the measurements are presented as a function of the layer thickness. Up to 2 orders of
15 14 13 12
Hardness (GPa)
11 10 9 8
Load untreated LEI PIII
20 mN 50 mN
Temperature 300 - 400 ˚C LEI 15 min, PIII 60 min
7 6 5 4 3 2 1 0
0.00
1
1
0
Layer Thickness (µm) Fig. 1. Effective surface hardness as a function of layer thickness. Please note the logarithmic depth scale.
444
D. Manova et al. / Nucl. Instr. and Meth. in Phys. Res. B 257 (2007) 442–446
Fig. 2. Wear rate as a function of layer thickness. Please note the logarithmic depth scale.
magnitude lower wear rates are observed after nitriding. However, in contrast to the thickness dependent hardness, nearly no influence of the layer thickness on the wear rates at low loads is seen for both methods of implantations. At high loads and thin layers, a different breakthrough is observed, which is not subject of this paper. The low wear, independent of layer thickness is now investigated further. Fig. 3 depicts one-dimensional profiles across wear tracks after different total length on two samples, non-implanted and implanted. The thickness of the nitrided layer was about 1.5 lm, as obtained from secondary ion mass spectroscopy measurements. Using a larger ball (B 3 mm) at about the same contact pressure leads to quantifiable wear tracks in the wear tests. A strong decrease of the wear track depth for the same total wear
path compared to the control sample is observed for the nitrided sample. Additional analysing of the 3D profiles yields a wear rate decrease by about a factor of 200 after implantation. However only slight increase in the wear rate of the nitrided sample of about factor of 3 is detected after total wear path length corresponding to 5–10 lm, which is far beyond the nitrided layer. The profiles for both samples are geometrically similar with the width given by contact area, which increases with the total wear path. This can be seen especially very good for the implanted sample after shorter wear test. Redeposition of material during the wear test is clearly seen for the non-treated sample, while it is nearly missing for the nitrided sample, not only at the edges but also at the reversal points.
Fig. 3. One-dimensional wear track profiles after different wear paths (different time for wear test) for untreated and nitrogen implanted (10 kV, 350 C, 2 h) samples.
D. Manova et al. / Nucl. Instr. and Meth. in Phys. Res. B 257 (2007) 442–446
Fig. 4. SEM micrographs in the vicinity of a wear track from cross section of untreated (a) and nitrided sample (b).
Metallographic cross-sections were made before and after wear test on implanted and the as-received samples. Wear tracks for this investigation were manufactured with the WC ball due to the reason mentioned above. Microscopic viewgraphs of the area in the vicinity of the wear tracks, obtained from SEM, are presented in Fig. 4(a) non-implanted sample and (b) a sample after nitrogen implantation. The microstructure just below the wear track on the non-treated sample shows strongly deformed grains in the direction perpendicular to the applied load and indicates abrasive mode after the shear strength is reached. In contrast, the implanted sample shows a microstructure with a drastic change of grain shape in the vicinity of the wear track, as the individual grains are barely visible and the structure resembles highly deformed material with strongly elongated grains. Plan view images of the samples, after wear experiments performed with a 10 lm diamond ball, are shown in Fig. 5, as obtained from optical microscope. Formation of a large pile-up and a huge number of dislocations is clearly visible in Fig. 5(a) presenting the microstructure of the nonimplanted sample. In contrast, the observable pile-up is very small and no visible dislocations can be found near the wear track for the implanted material (Fig. 5(b)).
445
Fig. 5. Plan view optical micrographs of a wear track on untreated (a) and nitrided (b) sample.
4. Discussion The presented results can be used to highlight the different deformation mechanism for the base material and for nitrided surfaces. For the non-implanted sample the applied load results in an agglomeration of dislocations leading to grain refinement with much smaller gain size in the area where the local stress exceeds the strength of stainless steel. The local pressure just below the contact area exceeding 1.0 GPa has to be put in relation to the strength of about 0.5 GPa [12]. A significant pile-up after wear tests on untreated austenitic stainless steel, formed by particles removed from the wear track, indicates abrasive wear correlated with ploughing into the surface by the counterbody ball. In contrast, the same applied load on the nitrided samples leads only to strong deformation of grains resulting in highly anisotropic thin fibres with less pronounced dislocations. The nearly missing pile-up for the nitrided sample, together with a non-existent weight gain of the WC balls, indicates a completely different wear mechanism. As the penetration depth of the ball at the same contact pressure is smaller for a higher hardness, respective elastic modulus of the surface, a transition from surface removal
446
D. Manova et al. / Nucl. Instr. and Meth. in Phys. Res. B 257 (2007) 442–446
to deformation due to a indentation below the threshold where cutting occurs can be envisaged [13]. For shallower cutting depths and the same tool geometry, only substantial subsurface plastic flow will occur instead of chip formation. However, an additional effect is proposed to transform the surface properties to explain the complete observations, including the missing dislocations within the nitrided surfaces. Cracking and abrasion are indicative of large intergranular stresses, in contrast to elastic intragranular stored energy [14], which can be correlated with the stacking fault energies [15]. Similar to high nitrogen steels containing molybdenum, a decreased stacking fault energy there results in a decreased fatigue crack growth rate there [16]. As only 0.2 wt% nitrogen were necessary to observe this effect, an even more pronounced manifestation can be expected at around 15–25 at% in the present case. As a consequence, the nitrided surface layer is not removed by the wear process but recycled and folded into the subsurface below the wear track during the actual wear process – similar to cold work hardening but more effective – where an additional wear resistant layer is formed dynamically during the wear process. This leads to an effective wear resistant layer about 3–5 times thicker than implanted zone.
5. Summary and conclusions Very hard and wear resistant layers are formed after nitrogen insertion into austenitic stainless steel using PIII and LEI. An outstanding wear resistance was observed for an effective layer thickness of about 3–5 times the thickness of the nitrided layer. A change from abrasive wear towards substantial subsurface plastic flow is mediated by the higher elastic modulus and the modified stacking fault energy of the expanded austenite. For applications this means a wear protection of the nitrided surface extends beyond the original layer thickness without thermal or
stress induced experiments.
nitrogen
diffusion
during
the
wear
Acknowledgements Parts of this work were supported by European Regional Development Fund (ERDF) and the State of Saxony (Project 106967/1650), as well as the German Federation of Industrial Research Associations (AiF, Project KF 0189601UK5). References [1] C. Blawert, B.L. Mordike, Y. Jira´skova´, O. Schneeweiss, Surf. Coat. Technol. 116–119 (1999) 189. [2] D.L. Williamson, J.A. Davis, P.J. Wilbur, J.J. Vajo, R. Wei, J.N. Matossian, Nucl. Instr. and Meth. B 127–128 (1997) 930. [3] G.A. Collins, R. Hutchings, K.T. Short, J. Tendys, Surf. Coat. Technol. 103–104 (1998) 212. [4] D. Manova, I.-M. Eichentopf, D. Hirsch, S. Ma¨ndl, H. Neumann, B. Rauschenbach, IEEE Trans. Plasma Sci. 34 (2006) 1136. [5] W.P. Tong, N.R. Tao, Z.B. Wang, J. Lu, K. Lu, Science 299 (2003) 686. [6] H. Bu¨ckle, Mikroha¨rtepru¨fung und ihre Anwendung, Berliner Union Verlag, Stuttgart, 1957. [7] M. Samandi, B.A. Shedden, D.I. Smith, G.A. Collins, R. Hutchings, J. Tendys, Surf. Coat. Technol. 59 (1993) 261. [8] E. Menthe, K.-T. Rie, Surf. Coat. Technol. 116–119 (1999) 199. [9] S. Ma¨ndl, R. Gu¨nzel, E. Richter, W. Mo¨ller, Surf. Coat. Technol. 100–101 (1998) 372. [10] D. Manova, S. Ma¨ndl, H. Neumann, B. Rauschenbach, Surf. Coat. Technol. 200 (2005) 137. [11] D. Manova, S. Ma¨ndl, B. Rauschenbach, Plasma Sources Sci. Technol. 10 (2001) 423. [12] L. Vitos, P.A. Korzhavyi, B. Johansson, Adv. Eng. Mater. 4 (2004) 228. [13] J.D. Thiele, S.N. Melkote, J. Manuf. Proc. 2 (2000) 270. [14] Y.-D. Wang, H. Tian, A.D. Stoica, X.-L. Wang, P.K. Liaw, J.W. Richardson, Nat. Mater. 2 (2003) 101. [15] A.G. Froseth, P.M. Derlet, H. Van Swygenhoven, Adv. Eng. Mater. 7 (2005) 16. [16] M. Murayama, K. Hono, H. Hirukawa, T. Ohmura, S. Matsuoka, Scripta Mater. 41 (1999) 467.