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Properties of duplex coatings prepared by plasma nitriding and PVD Ti–C:H deposition on X20Cr13 ferritic stainless steel
Thin Solid Films 322 Ž1998. 206–212
Properties of duplex coatings prepared by plasma nitriding and PVD Ti–C:H deposition on X20Cr13 ferritic stainless steel T. Michler ) , M. Grischke, K. Bewilogua, H. Dimigen Fraunhofer-Institut fur Bienroder Weg 54 E, 38108 Braunschweig, Germany ¨ Schicht- und Oberflachentechnik, ¨ Received 5 August 1997; accepted 16 October 1997
Abstract Duplex-coating procedures consisting of plasma nitriding and Me–C:H hard coating lead to an improved performance of the devices because the Me–C:H coating is supported by the nitrided phase and, therefore, the ‘eggshell-effect’ is avoided. Furthermore, this support leads to a higher load-bearing capacity of the thin film. Two standard procedures Žclassical high-pressure plasma nitriding and unbalanced magnetron sputtering of Ti–C:H. were performed subsequently to prepare the duplex coatings on X20Cr13 ferritic stainless steel. The corrosion resistance of the steel could be improved by nitriding at 4508C compared to the untreated ferritic substrate. The roughness is determined by the nitriding step. The weakest point of the coating is the transition zone between the nitrided and the untreated substrate and not the interface between the Ti–C:H coating and the nitrided substrate as shown by the Rockwell and scratch tests. The adhesion of the duplex structure Žplasma nitriding and Ti–C:H coating. on X20Cr13 Ž230 HV0.1. is superiour to the adhesion of Ti–C:H on hardened 100Cr6 ball-bearing steel Ž62 HRCf 820 HV0.1, not nitrided.. q 1998 Elsevier Science S.A. All rights reserved. Keywords: Coatings; Plasma nitriding; Steel
1. Introduction For many applications the performance of a coating is limited by the mechanical properties of the substrate material itself. Thin films require a certain support by the substrate material to avoid the ‘eggshell-effect’ and thus to obtain a good adhesion of the hard coating. Furthermore, this support leads to a higher load-bearing capacity of the thin film. Several authors recently reported on the combination of plasma nitriding and PVD hard coating of TiN or CrN either by a discontinuous process Žplasma nitriding and PVD in two runs. w1–6x or a continuous process, where nitriding and hard coating is done in the same run w5–9x. Considering the productivity of such a treatment, the latter is obviously more promising because process temperatures of both plasma nitriding and hard coating of, e.g., TiN are in the same range and the vessel has to be evacuated only once.
) Corresponding author. Tel.: q49-531-2155-601; fax: q49-5312155-904; e-mail: [email protected].
0040-6090r98r$19.00 q 1998 Elsevier Science S.A. All rights reserved. PII S 0 0 4 0 - 6 0 9 0 Ž 9 7 . 0 0 9 5 9 - 0
The optimisation of duplex coatings for construction and hot working steels was investigated by Dingremont et al. w3,6,7x. When choosing inadequate parameters, the formation of a ‘black layer’ at the TiN compound layer interface deteriorates the good properties of the duplex coating. The best results were reached by a continuous process. Improved wear resistance for metal injection mouldings was reported. Bucken et al. w8x and Hock ¨ ¨ et al. w9x reported on a significant increase of the critical load of TiN on prenitrided X155CrVMo121 cold working steel and S6-5-2 tool steel. Meletis et al. w2x reported on the combination of plasma nitriding and ion-beam deposition of DLC Ždiamond-like carbon. which improved the coating support and therefore improved the load-bearing capacity tested by a ball-on-disc test. A report about the combination of plasma nitriding and Me–C:H Žmetal containing amorphous hydrogenated carbon. hard coating could not be found. Metal–carbon films offer the combination of low coefficients of friction Ž m f 0.2 vs. steel. and a high abrasive as well as a high adhesive wear resistance w10–13x. Therefore they are very promising as engineering coatings for high-precision components. Applications are machine parts
T. Michler et al.r Thin Solid Films 322 (1998) 206–212
207
Fig. 1. Schematic structure of duplex films Žplasma nitriding plus hard coating..
like bearings, shafts or gears w10,14–16x, forming tools w16–18x, thread cutting tools w19x or valves w16x as well as special tribological applications in the automotive industry w20x. X20Cr13 stainless steel ŽGerman standard no. 1.4021, US: AISI420, France: Z20C13, UK: B.S.420S37. is a relatively cheap steel which is used in a lot of applications like shafts, bolts, pins and die-casting forms Žhardened material with good mechanical stability. or any kind of valve Žnormalized material with a good corrosion resistance.. Most of these applications are also typical for the use of Me–C:H coatings. Plasma nitriding of X20Cr13 at temperatures below 5008C offers the possibility of hardening the material while maintaining its corrosion resistance. This paper reports on some properties of duplex coatings Žplasma nitriding plus PVD Ti–C:H. on X20Cr13 ferritic stainless steel Žnormalized to 230 HV..
2. Experimental details Plasma nitriding of X20Cr13 ferritic stainless steel was performed in a PlaTeG-PP40 plasma nitriding plant Žpulsed DC. described elsewhere w21x. The pressure was set to 350
Pa. Nitriding time, N2rH 2 gas-flow ratio and bias voltage were varied in order to get different diffusion profiles. Plasma nitriding was performed at the following parameters: N2rH 2 gas-flow ratio: 0.09–4, nitriding time: 120– 360 min, bias voltage: 340–470 V. To avoid a reduction of the corrosion resistance of the steel, a low-temperature heat treatment at 4508C was chosen. The corrosion resistance was tested by the ferrirferrocyanide test Ž3 g NaCl q 0.1 g K 3 FeŽCN. 6 in 100 ml distilled water.. The roughness was measured mechanically using a profilometer ŽRank Taylor Hobson, UK.. The surface layer was analysed by ŽX-ray diffraction. XRD ŽCu K a radiation, Siemens, Germany. and the Young’s moduli were measured by nanoindentation ŽPhilips, Netherlands.. After nitriding, the samples were coated with a titanium-containing amorphous hydrogenated carbon film ŽTi–C:H. prepared by unbalanced magnetron sputtering Žf 1 Pa, C 2 H 2 , Ti-Target, - 2008C. in a TRITEC 1000 batch coater ŽLeybold, Germany. described elsewhere w22x. The thickness of the films varied between 2.7 and 3.7 m m for the films discussed here. Fig. 1 shows the schematic structure of the duplex structures thus obtained. The adhesion was tested by the Rockwell test ŽWolpert, Germany. and the scratch test ŽFraunhofer-IPA, Germany.,
Table 1 Surface hardness Sample No.
N2 r H 2 gas flow ratio
1 2 3 4 5
0.09 0.19 0.25 4
Nitriding time [min] not nitrided 240 180 240 240
Bias voltage [V]
Metallurgical phases (XRD)
Surface hardness [HV0.05]
460 460 420 420
ferrite not measured X ferrite, g ´ ´
230 1120 1070 1140 1160
T. Michler et al.r Thin Solid Films 322 (1998) 206–212
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Table 3 Corrosion resistance Metallurgical phases (XRD)
Corrosion resistance (Ferrir rferro-cyanide test)
ferrite X ferrite, g ´
good good excellent
The untreated material showed a poorer corrosion resistance compared to the plasma-nitrided steel with monophase Ž ´ .. An improvement in corrosion resistance could not be observed for the plasma-nitrided steel with multiphase Žferrite and g X , see Table 3. compared to the untreated material. When nitriding high-chromium steels at temperatures above 5008C, the corrosion resistance deteriorates because of the formation of chromium nitride w23,24x. This can be avoided by nitriding at lower temperatures. When nitriding X20Cr13 at 4508C, no formation of chromium nitride could be detected by XRD and the corrosion resistance even increased because iron nitrides have a better corrosion resistance than ferrite Ž a-Fe. w23x. Detailed results about the structure, corrosion resistance and topography of plasma-nitrided X20Cr13 will be published separately.
Fig. 2. Typical hardness–depth profile of plasma-nitrided X20Cr13.
respectively. The Rockwell test was performed according to the VDI-RL 3198 and the scratch test according to DIN V ENV 1071.
3. Results and discussion 3.1. Plasma nitriding of X20Cr13 As can be seen from Table 1 even at soft nitriding conditions Žlow N2rH 2 gas-flow ratio, short nitriding time or low bias voltage. a high surface hardness of about 1100 HV was reached. XRD measurements indicate that this is due to the formation of the g-phase Žsample 2, 3.. When the N2rH 2 gas-flow ratio exceeds 0.25 at these conditions, the formation of ´-phase could be observed and the surface hardness increased slightly to about 1200 HV Žsample 4, 5.. In our case, the ´-phase obviously consists of ŽFe,Cr. 2 – 3 N. Hardness–depth profiles performed at metallurgical cross sections showed that in all cases the nitriding depth was in the range of 15–25 m m. After that, the hardness decreased rapidly from 1100–1200 HV to the core hardness of 230 HV. Fig. 2 shows a typical hardness–depth profile. The surface roughness of the polished samples Ž R z,ISO f 0.1 m m. increased to R z,ISO ) 0.75 m m even at soft nitriding conditions Žsee Table 2.. This may limit the use of such duplex coatings and has to be kept in mind. Furthermore, it can be seen that the roughness is determined by the nitriding step because no significant increase was measured after the PVD-coating step.
3.2. Adhesion of duplex coatings on X20Cr13 When nitrided steels with a compound layer Žg X or ´ . are coated by PVD, inadequate PVD parameters can lead to a denitriding of the compound layer due to retrodiffusion of nitrogen which results in a bad adhesion of the PVD coating. This decomposed phase is called ‘black layer’ because it appears black in etched metallurgical cross sections. Investigations by Dingremont et al. w3x showed that a black layer occurs when the nitrided substrate is heated up too high before the PVD-coating step. Of course this temperature depends on the steel grade but it can be as low as 3508C in an argon plasma atmosphere. Our investigations showed no formation of a black layer after the deposition of Ti–C:H. This is obviously due to the low deposition temperature of carbon-based coatings which is usually below 2008C. At this temperature a retrodiffusion of nitrogen is very unlikely.
Table 2 Surface roughness of polished samples Sample No.
N2 r H 2 gas flow ratio
1 6 7 8
0.15 4 4
Nitriding time [min] not nitrided 240 120 360
Bias voltage [V]
R z,ISO after plasmanitriding [ m m]
460 420 420
0.81 0.81 1.02
R z,ISO after coating [ m m]
R z,ISO f 0.1 0.76 0.91 1.24
T. Michler et al.r Thin Solid Films 322 (1998) 206–212
209
The adhesion tests performed here have to be compared to those on 100Cr6 ball-bearing steel ŽDIN 1.3505, hardened to 62 HRCf 750 HV. which is usually used as a reference material. The properties of Ti–C:H films on polished 100Cr6 Ž R z,ISO s 0.04–0.07. are as follows: atomic ratio TirC ŽEPMA., f 0.25, film thickness, 2.7– 3.7 m m; Rockwell test adhesion class Žaccording to VDI-RL 3198., 1–2; critical load Lc Žfirst spallations., 35–40 N. 3.2.1. Rockwell test Rockwell tests of Ti–C:H on 100Cr6 steel usually show radial cracks and little cohesive spallations around the indentation. When evaluating the adhesion of a PVD coating on a plasma-nitrided substrate, it has to be taken into account that the weakest point of the duplex structure is not necessarily the interface between the hard coating and the substrate but may be located within the nitriding layer. Fig. 3 shows the typical Rockwell indentations of Ža. a Ti–C:H coated, Žb. a plasma-nitrided and Žc. a duplex treated Žplasma-nitrided plus Ti–C:H coated. sample. As can be seen, the adhesion of the Ti–C:H coating on the untreated substrate Ža. was very poor. This is due to the lack of support of the soft substrate material Ž230 HV.. Almost no difference could be detected between the indentations on the plasma-nitrided Žb. and the duplex-treated Žc. surface. In both cases widespread cracks are visible. In all cases examined here no spallation of the Ti–C:H film occurs when deposited on nitrided substrates. It has to be pointed out that the fractures around the indentation areas of the prenitrided surfaces were independent of the plasma-nitriding conditions. Even at soft nitriding conditions, this type of fracture occurred. To identify the failure mechanism in detail, metallurgical cross sections of the Rockwell indentations were made. Fig. 4 shows a cross section of a Rockwell indentation on a duplex-treated sample. It can be seen that the distance between the surface and the origin of the fractures is about 20 m m, which corresponds with the sharp drop in hardness between the plasma-nitrided and the untreated material Žsee Fig. 2.. This indicates that the weakest point of the duplex structure is the transition zone between the nitrided and the untreated substrate and not the interface between the Ti–C:H coating and the nitrided substrate. 3.2.2. Scratch test At moderate loads Žf 10–35 N. Me–C:H films generally show strain cracks behind the indenter due to the small fracture strains of hydrocarbon films w25x. At higher loads small spallations outside the track occur. This load is usually marked as ‘critical load Ž Lc .’. At higher loads the films are peeled off the substrate inside and outside the scratch track. Furthermore, it is well known that the critical load increases with increasing film thickness w26x. In all cases examined here, an identical scratch-test behaviour of all duplex-treated samples was observed. At moderate loads the coatings showed strain cracks within
Fig. 3. Typical Rockwell indentations of Ža. Ti–C:H coated, Žb. plasmanitrided and Žc. duplex-treated X20Cr13.
the track which were followed by wide shell-like spallations outside the track ŽFig. 5.. The first occurrence of these spallations was marked as the critical load Lc . It should be pointed out that no separate peeling of the Ti–C:H film from the nitrided surface could be observed during the whole scratch Žnormal load 0–100 N..
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T. Michler et al.r Thin Solid Films 322 (1998) 206–212
Fig. 6. Critical load of duplex coatings Žplasma nitriding plus Ti–C:H. on X20Cr13 depending on the nitriding time. Fig. 4. Cross section of a Rockwell indentation on duplex-treated X20Cr13.
Fig. 6 shows that the critical load in the scratch test increased almost linearly with the square root of nitriding time Ž t 0.5 . from less than 10 to 100 N. A nitriding time of only 120 min Ž t 0.5 f 11 min0.5 . leads to critical loads greater than 40 N, which is superiour to Ti–C:H on hardened 100Cr6. The nitriding time Ž t . is correlated to the nitriding depth Ž x . via the well-known t 0.5 law Žsee, e.g., Ref. w27x. x A 't . Because of the excellent adherence of the Ti–C:H film on the nitrided substrate, it can be assumed that Lc is proportional to the nitriding depth in the parameter field examined here and that Lc is dominated by the nitriding step. Critical loads measured on plasma-nitrided samples without a Ti–C:H coating showed the same values Žwithin the deviations. compared to plasma-nitrided samples with a Ti–C:H coating.
Fig. 5. Scratch track and ‘shell-like’ spallations on duplex-treated X20Cr13.
In Fig. 7 the critical load is plotted vs. the N2rH 2 gas-flow ratio at different substrate biases during nitriding. It can be seen in general that with increasing N2rH 2 gas-flow ratio the critical load increases Žup to about 80 N.. This is due to the higher concentration of active nitrogen. In detail, the critical loads increase rapidly above 60 N for N2rH 2 s 0 . . . 1 until they reach a saturation around 80 N at N2rH 2 G 1. Also here, the critical loads increase with increasing nitriding time Ž180 ™ 240 min. and a comparison with Fig. 6 indicates that a further increase in critical load can only be reached at longer nitriding times. Fig. 7 also shows that an increasing bias voltage overcompensates for a shorter nitriding time Ž Lc s 45 N at N2rH 2 s 0.25, 240 min, 420 V; L c s 71 N at N2rH 2 s 0.25, 180 min, 460 V.. This indicates that the bias voltage also influences the nitriding depth significantly. This can be understood if even at low N2rH 2 gas-flow ratios the concentration of active nitrogen can be increased with an increasing bias. These results indicate that critical loads up to 80 N can be reached by short-time nitriding ŽF 180 min. with optimized gas flows ŽN2rH 2 G 1. and substrate biases of 460 V or higher, which lead to a relatively cheap process due
Fig. 7. Critical load of duplex coatings Žplasma nitriding plus Ti–C:H. on X20Cr13 depending on the N2 rH 2 gas-flow ratio and the bias voltage of the nitriding process.
T. Michler et al.r Thin Solid Films 322 (1998) 206–212 Table 4 Mechanical properties Metallurgical phases (XRD)
Hardness [HV0.05]
Young’s modulus [GPa]
X20Cr13 Žuntreated material. X pn-X20Cr13; ferrite, g pn-X20Cr13; ´ Ti–C:H, atomic ratio TirC f 0.25
230 f1100 f1200 f1400
218"24 238"21 244"44 f130
to reduced nitriding times. Much higher bias voltage values may lead to an extreme rounding of edges and also to higher surface roughnesses. Interpreting Fig. 5, it can be seen that the origin of the shell-like spallations was within the substrate material which indicates that here also the weakest point of the duplex structure is not the interface between the Ti–C:H film and the substrate. Because of the excellent adherence of the Ti–C:H film on the nitrided substrate, it can be assumed that the high critical loads are due to the high thickness of the duplex structure Ž3 m m Ti–C:H plus 20–30 m m plasma-nitrided zone.. The mechanical properties of the materials which are of interest are listed in Table 4. It can be seen that the Young’s modulus of the nitrided phases are slightly higher compared to the untreated material. The high standard deviations are obviously due to the rough surface. However, the Ti–C:H film Ž E f 130 GPa. has a higher elasticity compared to the steel Žnitrided phases: E f 240 GPa, untreated material: E f 220 GPa.. The hardnesses of the Ti–C:H film Žf 1400 HV. and the nitrided phases Ž1100– 1200 HV. are about the same but there is a sharp drop in hardness in the plasma-nitrided and the untreated material Ž230 HV.. This indicates that the hardness of the substrate material is the dominating factor when a high load-bearing capacity of the functional hard coating is required. This can be reached either by using very hard substrate bulk materials or by applying only a supportive layer by plasma nitriding. Our results verify the results obtained by Oppel and Kayser w5x who reported on a significant increase in the critical load of TiN on X35CrMo17 stainless steel which was prenitrided at short times Ž- 180 min.. Accompanying investigations were performed on samples grinded to certain roughnesses Ž R z,ISO s 0.1–0.7 m m.. The critical load was not dependent on the surface roughness in this range.
4. Summary and conclusions Duplex coatings prepared by classical high-pressure plasma nitriding and PVD Ti–C:H hard coating on X20Cr13 ferritic stainless steel were investigated by means of the Rockwell test and scratch test.
211
It was found that the critical load of Ti–C:H films on X20Cr13 Ž230 HV. can be increased significantly by short-time nitriding Ž180 min. up to about 80 N. The critical loads of the duplex structure Žplasma nitriding and Ti–C:H coating. on X20Cr13 Ž230 HV0.1. are superiour to the critical loads of Ti–C:H on hardened 100Cr6 Ž820 HV0.1. ball-bearing steel Žnot nitrided.. The weakest point of the duplex structure on X20Cr13 is not the interface between the Ti–C:H coating and the nitrided substrate but the transition zone between the nitrided and the untreated substrate. Because of the low deposition temperatures of carbon-based coatings Žbelow 2008C. black-layer formation does not occur. The results obtained by Oppel and Kayser w5x show that this duplex structure also works with other high-chromium steels like X35CrMo17. This drastically widens the range of application for these kinds of steels. Acknowledgements The authors would like to thank C. Beckmann for preparing the metallurgical cross sections, K. Taube for making the nanoindentation measurements, C. Siebert for making the roughness measurements and A. Bulak from the Institut fur und Plasmatechnische ¨ Oberflachentechnik ¨ Werkstoffentwicklung, TU Braunschweig for making the XRD scans. This work was sponsored by the German Federal Ministry for Research and Technology Žbmb q f. under grant no. 13N6466. References w1x S.-C. Lee, W.-Y. Ho, W.-L. Pao, Surf. Coat. Technol. 73 Ž1995. 34–38. w2x E.I. Meletis, A. Erdemir, G.R. Fenske, Surf. Coat. Technol. 73 Ž1995. 39–45. w3x N. Dingremont, A. Parinelli, E. Bergmann, H. Michel, Surf. Coat. Technol. 61 Ž1993. 187–193. w4x D’Haen, C. Quaeyhaegens, L.M. Stals, Surf. Coat. Technol. 61 Ž1993. 194–200. w5x W. Oppel, O. Kayser, JOT 11 Ž1995. XVI–IXX. w6x N. Dingremont, E. Bergmann, P. Collignon, H. Michel, Surf. Coat. Technol. 72 Ž1995. 163–168. w7x N. Dingremont, E. Bergmann, P. Collignon, Surf. Coat. Technol. 72 Ž1995. 157–162. w8x B. Bucken, G. Leonhardt, R. Wilberg, K. Hock, ¨ ¨ H.-J. Spies, Surf. Coat. Technol. 68–69 Ž1995. 244–248. w9x K. Hock, H.-J. Spies, B. Larisch, Surf. ¨ G. Leonhardt, B. Bucken, ¨ Coat. Technol. 74–75 Ž1995. 339–344. w10x H. Dimigen, H. Hubsch, Philips Technol. Rev. 41 Ž6. Ž1983r1984. ¨ 186–197. w11x H. Dimigen, C.-P. Klages, Surf. Coat. Technol. 49 Ž1991. 543–547. w12x C.-P. Klages, R. Memming, Mater. Sci. For. 52–53 Ž1989. 609–644. w13x H. Dimigen, H. Hubsch, R. Memming, Appl. Phys. Lett. 50 Ž16. ¨ Ž1987. 1056–1058. w14x T. Lunow, R. Kocis, G. Leonhardt, R. Wilberg, Surf. Coat. Technol. 76–77 Ž1995. 579–582. w15x M. Grischke, JOT 1 Ž1994. 40–43. w16x A. Matthews, S.S. Eskildsen, Diamond Rel. Mater. 3 Ž1994. 902– 911.
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w17x K. Taube, to be published in Proc. PSE96. w18x K. Taube, M. Grischke, K. Bewilogua, Surf. Coat. Technol. 68–69 Ž1994. 662–668. w19x Contact Ž5., 1996, customs information brochure of Balzers, Lichtenstein. w20x J. Guttler, J. Reschke, Surf. Coat. Technol. 60 Ž1993. 531–535. ¨ w21x R. Grun, Mater. Sci. Eng. A 140 Ž1991. 435–441. ¨ H.-J. Gunther, ¨ w22x D. Hofmann, H. Schuessler, K. Bewilogua, H. Hubsch, J. Lemke, ¨ Surf. Coat. Technol. 73 Ž1995. 137–141.
w23x J. Elwart, Z. wirtschaftliche Fertig. 73 Ž1978. 12–18. w24x N. Dingremont, E. Bergmann, M. Hans, P. Collignon, Surf. Coat. Technol. 76–77 Ž1995. 218–224. w25x M. Fryda, C. Benndorf, C.-P. Klages, K. Taube, Diamond Rel. Mater. 1 Ž1992. 558–562. w26x M. Grischke, K. Bewilogua, K. Trojan, H. Dimigen, Surf. Coat. Technol. 74–75 Ž1995. 739–745. w27x W. Schatt ŽEd.., Einfuhrung in die Werkstoffwissenschaft, 6th edn. ¨ Leipzig: Dt. Verl. f. Grundstoffind., 1983.
Properties of duplex coatings prepared by plasma nitriding and PVD Ti–C:H deposition on X20Cr13 ferritic stainless steel T. Michler ) , M. Grischke, K. Bewilogua, H. Dimigen Fraunhofer-Institut fur Bienroder Weg 54 E, 38108 Braunschweig, Germany ¨ Schicht- und Oberflachentechnik, ¨ Received 5 August 1997; accepted 16 October 1997
Abstract Duplex-coating procedures consisting of plasma nitriding and Me–C:H hard coating lead to an improved performance of the devices because the Me–C:H coating is supported by the nitrided phase and, therefore, the ‘eggshell-effect’ is avoided. Furthermore, this support leads to a higher load-bearing capacity of the thin film. Two standard procedures Žclassical high-pressure plasma nitriding and unbalanced magnetron sputtering of Ti–C:H. were performed subsequently to prepare the duplex coatings on X20Cr13 ferritic stainless steel. The corrosion resistance of the steel could be improved by nitriding at 4508C compared to the untreated ferritic substrate. The roughness is determined by the nitriding step. The weakest point of the coating is the transition zone between the nitrided and the untreated substrate and not the interface between the Ti–C:H coating and the nitrided substrate as shown by the Rockwell and scratch tests. The adhesion of the duplex structure Žplasma nitriding and Ti–C:H coating. on X20Cr13 Ž230 HV0.1. is superiour to the adhesion of Ti–C:H on hardened 100Cr6 ball-bearing steel Ž62 HRCf 820 HV0.1, not nitrided.. q 1998 Elsevier Science S.A. All rights reserved. Keywords: Coatings; Plasma nitriding; Steel
1. Introduction For many applications the performance of a coating is limited by the mechanical properties of the substrate material itself. Thin films require a certain support by the substrate material to avoid the ‘eggshell-effect’ and thus to obtain a good adhesion of the hard coating. Furthermore, this support leads to a higher load-bearing capacity of the thin film. Several authors recently reported on the combination of plasma nitriding and PVD hard coating of TiN or CrN either by a discontinuous process Žplasma nitriding and PVD in two runs. w1–6x or a continuous process, where nitriding and hard coating is done in the same run w5–9x. Considering the productivity of such a treatment, the latter is obviously more promising because process temperatures of both plasma nitriding and hard coating of, e.g., TiN are in the same range and the vessel has to be evacuated only once.
) Corresponding author. Tel.: q49-531-2155-601; fax: q49-5312155-904; e-mail: [email protected].
0040-6090r98r$19.00 q 1998 Elsevier Science S.A. All rights reserved. PII S 0 0 4 0 - 6 0 9 0 Ž 9 7 . 0 0 9 5 9 - 0
The optimisation of duplex coatings for construction and hot working steels was investigated by Dingremont et al. w3,6,7x. When choosing inadequate parameters, the formation of a ‘black layer’ at the TiN compound layer interface deteriorates the good properties of the duplex coating. The best results were reached by a continuous process. Improved wear resistance for metal injection mouldings was reported. Bucken et al. w8x and Hock ¨ ¨ et al. w9x reported on a significant increase of the critical load of TiN on prenitrided X155CrVMo121 cold working steel and S6-5-2 tool steel. Meletis et al. w2x reported on the combination of plasma nitriding and ion-beam deposition of DLC Ždiamond-like carbon. which improved the coating support and therefore improved the load-bearing capacity tested by a ball-on-disc test. A report about the combination of plasma nitriding and Me–C:H Žmetal containing amorphous hydrogenated carbon. hard coating could not be found. Metal–carbon films offer the combination of low coefficients of friction Ž m f 0.2 vs. steel. and a high abrasive as well as a high adhesive wear resistance w10–13x. Therefore they are very promising as engineering coatings for high-precision components. Applications are machine parts
T. Michler et al.r Thin Solid Films 322 (1998) 206–212
207
Fig. 1. Schematic structure of duplex films Žplasma nitriding plus hard coating..
like bearings, shafts or gears w10,14–16x, forming tools w16–18x, thread cutting tools w19x or valves w16x as well as special tribological applications in the automotive industry w20x. X20Cr13 stainless steel ŽGerman standard no. 1.4021, US: AISI420, France: Z20C13, UK: B.S.420S37. is a relatively cheap steel which is used in a lot of applications like shafts, bolts, pins and die-casting forms Žhardened material with good mechanical stability. or any kind of valve Žnormalized material with a good corrosion resistance.. Most of these applications are also typical for the use of Me–C:H coatings. Plasma nitriding of X20Cr13 at temperatures below 5008C offers the possibility of hardening the material while maintaining its corrosion resistance. This paper reports on some properties of duplex coatings Žplasma nitriding plus PVD Ti–C:H. on X20Cr13 ferritic stainless steel Žnormalized to 230 HV..
2. Experimental details Plasma nitriding of X20Cr13 ferritic stainless steel was performed in a PlaTeG-PP40 plasma nitriding plant Žpulsed DC. described elsewhere w21x. The pressure was set to 350
Pa. Nitriding time, N2rH 2 gas-flow ratio and bias voltage were varied in order to get different diffusion profiles. Plasma nitriding was performed at the following parameters: N2rH 2 gas-flow ratio: 0.09–4, nitriding time: 120– 360 min, bias voltage: 340–470 V. To avoid a reduction of the corrosion resistance of the steel, a low-temperature heat treatment at 4508C was chosen. The corrosion resistance was tested by the ferrirferrocyanide test Ž3 g NaCl q 0.1 g K 3 FeŽCN. 6 in 100 ml distilled water.. The roughness was measured mechanically using a profilometer ŽRank Taylor Hobson, UK.. The surface layer was analysed by ŽX-ray diffraction. XRD ŽCu K a radiation, Siemens, Germany. and the Young’s moduli were measured by nanoindentation ŽPhilips, Netherlands.. After nitriding, the samples were coated with a titanium-containing amorphous hydrogenated carbon film ŽTi–C:H. prepared by unbalanced magnetron sputtering Žf 1 Pa, C 2 H 2 , Ti-Target, - 2008C. in a TRITEC 1000 batch coater ŽLeybold, Germany. described elsewhere w22x. The thickness of the films varied between 2.7 and 3.7 m m for the films discussed here. Fig. 1 shows the schematic structure of the duplex structures thus obtained. The adhesion was tested by the Rockwell test ŽWolpert, Germany. and the scratch test ŽFraunhofer-IPA, Germany.,
Table 1 Surface hardness Sample No.
N2 r H 2 gas flow ratio
1 2 3 4 5
0.09 0.19 0.25 4
Nitriding time [min] not nitrided 240 180 240 240
Bias voltage [V]
Metallurgical phases (XRD)
Surface hardness [HV0.05]
460 460 420 420
ferrite not measured X ferrite, g ´ ´
230 1120 1070 1140 1160
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Table 3 Corrosion resistance Metallurgical phases (XRD)
Corrosion resistance (Ferrir rferro-cyanide test)
ferrite X ferrite, g ´
good good excellent
The untreated material showed a poorer corrosion resistance compared to the plasma-nitrided steel with monophase Ž ´ .. An improvement in corrosion resistance could not be observed for the plasma-nitrided steel with multiphase Žferrite and g X , see Table 3. compared to the untreated material. When nitriding high-chromium steels at temperatures above 5008C, the corrosion resistance deteriorates because of the formation of chromium nitride w23,24x. This can be avoided by nitriding at lower temperatures. When nitriding X20Cr13 at 4508C, no formation of chromium nitride could be detected by XRD and the corrosion resistance even increased because iron nitrides have a better corrosion resistance than ferrite Ž a-Fe. w23x. Detailed results about the structure, corrosion resistance and topography of plasma-nitrided X20Cr13 will be published separately.
Fig. 2. Typical hardness–depth profile of plasma-nitrided X20Cr13.
respectively. The Rockwell test was performed according to the VDI-RL 3198 and the scratch test according to DIN V ENV 1071.
3. Results and discussion 3.1. Plasma nitriding of X20Cr13 As can be seen from Table 1 even at soft nitriding conditions Žlow N2rH 2 gas-flow ratio, short nitriding time or low bias voltage. a high surface hardness of about 1100 HV was reached. XRD measurements indicate that this is due to the formation of the g-phase Žsample 2, 3.. When the N2rH 2 gas-flow ratio exceeds 0.25 at these conditions, the formation of ´-phase could be observed and the surface hardness increased slightly to about 1200 HV Žsample 4, 5.. In our case, the ´-phase obviously consists of ŽFe,Cr. 2 – 3 N. Hardness–depth profiles performed at metallurgical cross sections showed that in all cases the nitriding depth was in the range of 15–25 m m. After that, the hardness decreased rapidly from 1100–1200 HV to the core hardness of 230 HV. Fig. 2 shows a typical hardness–depth profile. The surface roughness of the polished samples Ž R z,ISO f 0.1 m m. increased to R z,ISO ) 0.75 m m even at soft nitriding conditions Žsee Table 2.. This may limit the use of such duplex coatings and has to be kept in mind. Furthermore, it can be seen that the roughness is determined by the nitriding step because no significant increase was measured after the PVD-coating step.
3.2. Adhesion of duplex coatings on X20Cr13 When nitrided steels with a compound layer Žg X or ´ . are coated by PVD, inadequate PVD parameters can lead to a denitriding of the compound layer due to retrodiffusion of nitrogen which results in a bad adhesion of the PVD coating. This decomposed phase is called ‘black layer’ because it appears black in etched metallurgical cross sections. Investigations by Dingremont et al. w3x showed that a black layer occurs when the nitrided substrate is heated up too high before the PVD-coating step. Of course this temperature depends on the steel grade but it can be as low as 3508C in an argon plasma atmosphere. Our investigations showed no formation of a black layer after the deposition of Ti–C:H. This is obviously due to the low deposition temperature of carbon-based coatings which is usually below 2008C. At this temperature a retrodiffusion of nitrogen is very unlikely.
Table 2 Surface roughness of polished samples Sample No.
N2 r H 2 gas flow ratio
1 6 7 8
0.15 4 4
Nitriding time [min] not nitrided 240 120 360
Bias voltage [V]
R z,ISO after plasmanitriding [ m m]
460 420 420
0.81 0.81 1.02
R z,ISO after coating [ m m]
R z,ISO f 0.1 0.76 0.91 1.24
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The adhesion tests performed here have to be compared to those on 100Cr6 ball-bearing steel ŽDIN 1.3505, hardened to 62 HRCf 750 HV. which is usually used as a reference material. The properties of Ti–C:H films on polished 100Cr6 Ž R z,ISO s 0.04–0.07. are as follows: atomic ratio TirC ŽEPMA., f 0.25, film thickness, 2.7– 3.7 m m; Rockwell test adhesion class Žaccording to VDI-RL 3198., 1–2; critical load Lc Žfirst spallations., 35–40 N. 3.2.1. Rockwell test Rockwell tests of Ti–C:H on 100Cr6 steel usually show radial cracks and little cohesive spallations around the indentation. When evaluating the adhesion of a PVD coating on a plasma-nitrided substrate, it has to be taken into account that the weakest point of the duplex structure is not necessarily the interface between the hard coating and the substrate but may be located within the nitriding layer. Fig. 3 shows the typical Rockwell indentations of Ža. a Ti–C:H coated, Žb. a plasma-nitrided and Žc. a duplex treated Žplasma-nitrided plus Ti–C:H coated. sample. As can be seen, the adhesion of the Ti–C:H coating on the untreated substrate Ža. was very poor. This is due to the lack of support of the soft substrate material Ž230 HV.. Almost no difference could be detected between the indentations on the plasma-nitrided Žb. and the duplex-treated Žc. surface. In both cases widespread cracks are visible. In all cases examined here no spallation of the Ti–C:H film occurs when deposited on nitrided substrates. It has to be pointed out that the fractures around the indentation areas of the prenitrided surfaces were independent of the plasma-nitriding conditions. Even at soft nitriding conditions, this type of fracture occurred. To identify the failure mechanism in detail, metallurgical cross sections of the Rockwell indentations were made. Fig. 4 shows a cross section of a Rockwell indentation on a duplex-treated sample. It can be seen that the distance between the surface and the origin of the fractures is about 20 m m, which corresponds with the sharp drop in hardness between the plasma-nitrided and the untreated material Žsee Fig. 2.. This indicates that the weakest point of the duplex structure is the transition zone between the nitrided and the untreated substrate and not the interface between the Ti–C:H coating and the nitrided substrate. 3.2.2. Scratch test At moderate loads Žf 10–35 N. Me–C:H films generally show strain cracks behind the indenter due to the small fracture strains of hydrocarbon films w25x. At higher loads small spallations outside the track occur. This load is usually marked as ‘critical load Ž Lc .’. At higher loads the films are peeled off the substrate inside and outside the scratch track. Furthermore, it is well known that the critical load increases with increasing film thickness w26x. In all cases examined here, an identical scratch-test behaviour of all duplex-treated samples was observed. At moderate loads the coatings showed strain cracks within
Fig. 3. Typical Rockwell indentations of Ža. Ti–C:H coated, Žb. plasmanitrided and Žc. duplex-treated X20Cr13.
the track which were followed by wide shell-like spallations outside the track ŽFig. 5.. The first occurrence of these spallations was marked as the critical load Lc . It should be pointed out that no separate peeling of the Ti–C:H film from the nitrided surface could be observed during the whole scratch Žnormal load 0–100 N..
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T. Michler et al.r Thin Solid Films 322 (1998) 206–212
Fig. 6. Critical load of duplex coatings Žplasma nitriding plus Ti–C:H. on X20Cr13 depending on the nitriding time. Fig. 4. Cross section of a Rockwell indentation on duplex-treated X20Cr13.
Fig. 6 shows that the critical load in the scratch test increased almost linearly with the square root of nitriding time Ž t 0.5 . from less than 10 to 100 N. A nitriding time of only 120 min Ž t 0.5 f 11 min0.5 . leads to critical loads greater than 40 N, which is superiour to Ti–C:H on hardened 100Cr6. The nitriding time Ž t . is correlated to the nitriding depth Ž x . via the well-known t 0.5 law Žsee, e.g., Ref. w27x. x A 't . Because of the excellent adherence of the Ti–C:H film on the nitrided substrate, it can be assumed that Lc is proportional to the nitriding depth in the parameter field examined here and that Lc is dominated by the nitriding step. Critical loads measured on plasma-nitrided samples without a Ti–C:H coating showed the same values Žwithin the deviations. compared to plasma-nitrided samples with a Ti–C:H coating.
Fig. 5. Scratch track and ‘shell-like’ spallations on duplex-treated X20Cr13.
In Fig. 7 the critical load is plotted vs. the N2rH 2 gas-flow ratio at different substrate biases during nitriding. It can be seen in general that with increasing N2rH 2 gas-flow ratio the critical load increases Žup to about 80 N.. This is due to the higher concentration of active nitrogen. In detail, the critical loads increase rapidly above 60 N for N2rH 2 s 0 . . . 1 until they reach a saturation around 80 N at N2rH 2 G 1. Also here, the critical loads increase with increasing nitriding time Ž180 ™ 240 min. and a comparison with Fig. 6 indicates that a further increase in critical load can only be reached at longer nitriding times. Fig. 7 also shows that an increasing bias voltage overcompensates for a shorter nitriding time Ž Lc s 45 N at N2rH 2 s 0.25, 240 min, 420 V; L c s 71 N at N2rH 2 s 0.25, 180 min, 460 V.. This indicates that the bias voltage also influences the nitriding depth significantly. This can be understood if even at low N2rH 2 gas-flow ratios the concentration of active nitrogen can be increased with an increasing bias. These results indicate that critical loads up to 80 N can be reached by short-time nitriding ŽF 180 min. with optimized gas flows ŽN2rH 2 G 1. and substrate biases of 460 V or higher, which lead to a relatively cheap process due
Fig. 7. Critical load of duplex coatings Žplasma nitriding plus Ti–C:H. on X20Cr13 depending on the N2 rH 2 gas-flow ratio and the bias voltage of the nitriding process.
T. Michler et al.r Thin Solid Films 322 (1998) 206–212 Table 4 Mechanical properties Metallurgical phases (XRD)
Hardness [HV0.05]
Young’s modulus [GPa]
X20Cr13 Žuntreated material. X pn-X20Cr13; ferrite, g pn-X20Cr13; ´ Ti–C:H, atomic ratio TirC f 0.25
230 f1100 f1200 f1400
218"24 238"21 244"44 f130
to reduced nitriding times. Much higher bias voltage values may lead to an extreme rounding of edges and also to higher surface roughnesses. Interpreting Fig. 5, it can be seen that the origin of the shell-like spallations was within the substrate material which indicates that here also the weakest point of the duplex structure is not the interface between the Ti–C:H film and the substrate. Because of the excellent adherence of the Ti–C:H film on the nitrided substrate, it can be assumed that the high critical loads are due to the high thickness of the duplex structure Ž3 m m Ti–C:H plus 20–30 m m plasma-nitrided zone.. The mechanical properties of the materials which are of interest are listed in Table 4. It can be seen that the Young’s modulus of the nitrided phases are slightly higher compared to the untreated material. The high standard deviations are obviously due to the rough surface. However, the Ti–C:H film Ž E f 130 GPa. has a higher elasticity compared to the steel Žnitrided phases: E f 240 GPa, untreated material: E f 220 GPa.. The hardnesses of the Ti–C:H film Žf 1400 HV. and the nitrided phases Ž1100– 1200 HV. are about the same but there is a sharp drop in hardness in the plasma-nitrided and the untreated material Ž230 HV.. This indicates that the hardness of the substrate material is the dominating factor when a high load-bearing capacity of the functional hard coating is required. This can be reached either by using very hard substrate bulk materials or by applying only a supportive layer by plasma nitriding. Our results verify the results obtained by Oppel and Kayser w5x who reported on a significant increase in the critical load of TiN on X35CrMo17 stainless steel which was prenitrided at short times Ž- 180 min.. Accompanying investigations were performed on samples grinded to certain roughnesses Ž R z,ISO s 0.1–0.7 m m.. The critical load was not dependent on the surface roughness in this range.
4. Summary and conclusions Duplex coatings prepared by classical high-pressure plasma nitriding and PVD Ti–C:H hard coating on X20Cr13 ferritic stainless steel were investigated by means of the Rockwell test and scratch test.
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It was found that the critical load of Ti–C:H films on X20Cr13 Ž230 HV. can be increased significantly by short-time nitriding Ž180 min. up to about 80 N. The critical loads of the duplex structure Žplasma nitriding and Ti–C:H coating. on X20Cr13 Ž230 HV0.1. are superiour to the critical loads of Ti–C:H on hardened 100Cr6 Ž820 HV0.1. ball-bearing steel Žnot nitrided.. The weakest point of the duplex structure on X20Cr13 is not the interface between the Ti–C:H coating and the nitrided substrate but the transition zone between the nitrided and the untreated substrate. Because of the low deposition temperatures of carbon-based coatings Žbelow 2008C. black-layer formation does not occur. The results obtained by Oppel and Kayser w5x show that this duplex structure also works with other high-chromium steels like X35CrMo17. This drastically widens the range of application for these kinds of steels. Acknowledgements The authors would like to thank C. Beckmann for preparing the metallurgical cross sections, K. Taube for making the nanoindentation measurements, C. Siebert for making the roughness measurements and A. Bulak from the Institut fur und Plasmatechnische ¨ Oberflachentechnik ¨ Werkstoffentwicklung, TU Braunschweig for making the XRD scans. This work was sponsored by the German Federal Ministry for Research and Technology Žbmb q f. under grant no. 13N6466. References w1x S.-C. Lee, W.-Y. Ho, W.-L. Pao, Surf. Coat. Technol. 73 Ž1995. 34–38. w2x E.I. Meletis, A. Erdemir, G.R. Fenske, Surf. Coat. Technol. 73 Ž1995. 39–45. w3x N. Dingremont, A. Parinelli, E. Bergmann, H. Michel, Surf. Coat. Technol. 61 Ž1993. 187–193. w4x D’Haen, C. Quaeyhaegens, L.M. Stals, Surf. Coat. Technol. 61 Ž1993. 194–200. w5x W. Oppel, O. Kayser, JOT 11 Ž1995. XVI–IXX. w6x N. Dingremont, E. Bergmann, P. Collignon, H. Michel, Surf. Coat. Technol. 72 Ž1995. 163–168. w7x N. Dingremont, E. Bergmann, P. Collignon, Surf. Coat. Technol. 72 Ž1995. 157–162. w8x B. Bucken, G. Leonhardt, R. Wilberg, K. Hock, ¨ ¨ H.-J. Spies, Surf. Coat. Technol. 68–69 Ž1995. 244–248. w9x K. Hock, H.-J. Spies, B. Larisch, Surf. ¨ G. Leonhardt, B. Bucken, ¨ Coat. Technol. 74–75 Ž1995. 339–344. w10x H. Dimigen, H. Hubsch, Philips Technol. Rev. 41 Ž6. Ž1983r1984. ¨ 186–197. w11x H. Dimigen, C.-P. Klages, Surf. Coat. Technol. 49 Ž1991. 543–547. w12x C.-P. Klages, R. Memming, Mater. Sci. For. 52–53 Ž1989. 609–644. w13x H. Dimigen, H. Hubsch, R. Memming, Appl. Phys. Lett. 50 Ž16. ¨ Ž1987. 1056–1058. w14x T. Lunow, R. Kocis, G. Leonhardt, R. Wilberg, Surf. Coat. Technol. 76–77 Ž1995. 579–582. w15x M. Grischke, JOT 1 Ž1994. 40–43. w16x A. Matthews, S.S. Eskildsen, Diamond Rel. Mater. 3 Ž1994. 902– 911.
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