Corrosion–wear mechanisms of hard coated austenitic 316L stainless steels

Corrosion–wear mechanisms of hard coated austenitic 316L stainless steels

Wear 256 (2004) 491–499 Corrosion–wear mechanisms of hard coated austenitic 316L stainless steels Peter A. Dearnley∗ , Giles Aldrich-Smith1 School of...

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Wear 256 (2004) 491–499

Corrosion–wear mechanisms of hard coated austenitic 316L stainless steels Peter A. Dearnley∗ , Giles Aldrich-Smith1 School of Mechanical Engineering, University of Leeds, Leeds LS2 9JT, UK Received 9 August 2002; received in revised form 20 January 2003; accepted 30 January 2003

Abstract Austenitic stainless steels like 316L are amongst the most commonly selected structural alloys for use in corrosion environments. Unfortunately, their resistance to surface degradation caused during sliding contacts with other materials, in such environments is poor. Here, a synergistic combination of mechanical (wear) and chemical (corrosion) processes, known as corrosion–wear processes, are responsible for causing surface material loss. Accordingly, efforts are being made to identify surface treatments that can enhance the corrosion–wear resistance of 316L and similar alloys. One plausible solution is to apply thin hard coatings (∼5–10 ␮m thick) using various plasma-based technologies. In practice, this is often fraught with difficulty because of the complex nature of the pervading corrosion–wear mechanisms. This paper presents our recent work that has identified three major corrosion–wear mechanisms that must be minimised if a successful surface engineering design is to be achieved for corrosion–wear protection. These are: Type I—the removal of the coating passive film during sliding contact; Type II—galvanic attack of the substrate resulting in blistering of the coating and; Type III—galvanic attack of the counterface material leading to abrasion of the coating during subsequent sliding contact. © 2003 Elsevier B.V. All rights reserved. Keywords: S-phase; CrN; Corrosion–wear; Sputtering

1. Introduction Austenitic stainless steels are well-known for their corrosion resistance. However, when such materials are required to make sliding contact with other materials, in corrosion environments, a synergistic combination of mechanical (wear) and chemical (corrosion) processes, known as corrosion–wear processes, take place. Such processes have been investigated by many researchers and have been known for several decades [1,2]. The various synergies cause surface material loss that exceeds that attainable by either pure wear or pure corrosion processes alone. In regard to overcoming the problem of corrosion–wear of austenitic stainless steels, the simple approach of merely hardening the surface, through for example, conventional nitriding (>600 ◦ C), usually results in compromising of corrosion resistance [3]. Another approach is therefore required. ∗ Corresponding author. Tel.: +44-113-343-2141; fax: +44-113-233-2150. E-mail address: [email protected] (P.A. Dearnley). 1 Present address: The National Physical Laboratory, Teddington, Middlesex, UK.

0043-1648/$ – see front matter © 2003 Elsevier B.V. All rights reserved. doi:10.1016/S0043-1648(03)00559-3

S-phase is a meta-stable nitrogen supersaturated phase, based on Fe–Cr–Ni–N austenite, which shows considerable promise for improving corrosion–wear resistance, when it is applied as a coating, since it is both extremely hard and corrosion resistant. Although originally discovered after nitriding austenitic stainless steels at temperatures ∼400–500 ◦ C [3–5], we have shown that it is also possible to synthesise this cubic crystalline phase as a coating using reactive magnetron sputtering [6]. In some respects, this is a more flexible approach. Such coatings could be applied to a range of substrate materials. We have previously shown that magnetron sputtered S-phase coatings can be produced with nitrogen contents ranging from 7.9 ≤ [N] ≤ 32 at.% [7]. Such coatings increase linearly in hardness with increasing [N] reaching a maximum of ∼20 GPa (HV ∼ 2000 kg/mm2 ) when nitrogen concentrations are allowed to reach 28–30 at.% (Fig. 1). A measure of coating toughness is Lcrack , the load required to cause the onset of ring cracks, an effect that is common for S-phase coatings applied to 316L substrates [7]. Such tests show that S-phase coating toughness decreases with increasing [N], in the same range of [N] where hardness increases (Fig. 1). An alternative to S-phase coatings is to use a CrN coating. It was therefore of interest to compare the corrosion–wear

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Fig. 1. Influence of nitrogen content on the mechanical properties (hardness, elastic modulus and toughness) of nitrogen-alloyed stainless steel coatings (courtesy [7]).

behaviour of these materials, as well as to discover how the corrosion–wear response of S-phase coatings depend upon [N]. Through these investigations, it became possible to identify three important types of corrosion–wear mechanism that affect 316L when it is coated with hard layers of S-phase or CrN. The factors that determine the predominance of these mechanisms were also elucidated.

2. Experimental

compositions [7], [N] was determined. The three S-phase coating types, S-phase(1), S-phase(2) and S-phase(3) were, respectively, determined to contain 10.4, 25.4 and 23 at.% N. Coating thickness and microhardness data for the S-phase and CrN coatings are cited in Table 1. Microhardness of the coatings was determined using the well-known estimation method of Vingsbo et al. [8]. Here, a series of surface hardness data, determined using several hardness loads, is graphically plotted. By appropriate analysis, this then enables the influence of the lower substrate hardness to be systematically taken into account and allows the true coating hardness to be determined.

2.1. Coatings 2.2. Electrochemical scratch tests All the coatings (∼6–8 ␮m in thickness) were applied to polished 316L substrates using reactive unbalanced magnetron sputtering. Two experimental sputtered S-phase coatings, designated S-phase(1) and S-phase(2), were applied using a laboratory facility described in a previous publication [6]. One further experimental S-phase coating, designated S-phase(3) and a commercial CrN coating, were applied by Teer Coatings Ltd. (a UK coating vendor). The nitrogen concentration of the coatings [N] was determined indirectly. First, X-ray diffraction was used to determine the lattice parameter (a) of the S-phase unit cell for each of the coatings. Then, by using a calibration curve of lattice parameter versus [N], previously obtained for a range of S-phase

One form of corrosion–wear, described in Section 4 as “Type I” can be studied by using an electrochemical scratch test. In the Leeds design [9], a scratch indenter is made to reciprocate over the test surface, while monitoring the corrosion current and/or potential. In the preferred method, the value of the free corrosion potential (Ecorr ) is determined in situ by carrying out a narrow range anodic polarisation sweep. This potential is then applied to the test-piece and held constant, establishing a potentiostatic condition. Scratching is then commenced. The determined corrosion current (initially corresponding to Icorr ) rapidly rises when the passive layer is removed or damaged by scratching.

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Table 1 Basic coating data Coating material and designation

Nitrogen content (at.%)

Thickness (␮m)

Vickers microhardness (kg/mm2 )

Estimated yield strength (MPa)

Estimated shear yield strength, τ y (MPa)

Uncoated 316L S-phase(1) S-phase(2) S-phase(3) CrN

– 10.4 25.4 23.0 ∼50a

– 7 6 8 6

232 971 1598 1452 2263

759 3175 5225 4748 7400

379 1588 2613 2374 3700

a

Not quantified.

Table 2 Corrosion–wear testing regimes Test material

Test type

Ball type

Uncoated 316L

Reciprocation ball-on-plate Reciprocation ball-on-plate Reciprocation ball-on-plate Unidirectional ball-on-disc Unidirectional ball-on-disc Unidirectional ball-on-disc

Al2 O3

S-phase(1) coated 316L S-phase(2) coated 316L Uncoated 316L S-phase(3) coated 316L CrN coated 316L

Ball diameter (mm)

Test solution

Test load (N)

Estimated maximum Hertzian contact pressure P0 , (MPa)

Sliding velocity (m/s)

7.92

3% NaCl

1.47

716

0.03

Al2 O3

7.92

3% NaCl

1.47

806

0.03

Al2 O3

7.92

3% NaCl

1.47

826

0.03

WC–6% Co

10.0

3% NaCl and 6% FeCl3

1.76

648

0.24

WC–6% Co

10.0

3% NaCl and 6% FeCl3

1.76

742

0.24

WC–6% Co

10.0

3% NaCl and 6% FeCl3

1.76

782

0.24

Small oscillations in the corrosion current take place during the repetitive scratch phase of the test. Analysis of this behaviour shows that each small rise and fall in corrosion current is in phase with the cyclic motion of the scratch

Fig. 2. Average corrosion currents determined during electrochemical scratch tests carried out in 3% NaCl using an Al2 O3 indenter with a 200 g load. Error bars correspond to the maximum and minimum currents detected for each test. The data was collated over 800 s and corresponded to several hundred scratch cycles.

indenter. As the scratch indenter passes over the surface, the passive film is removed or damaged, but then begins to immediately reform. The restored and very thin film is short lived and becomes removed during the next scratch cycle. Upon ending the scratch cycle sequence, the passive film is able to grow to its equilibrium thickness and the corrosion current returns to its original and relatively low value (Icorr ). The mean of the corrosion current oscillations, during the scratching phase of such tests, referred to as the “average scratch current” within the confines of this paper, was found to vary for the various coated and uncoated test-pieces. Details are given in Section 3.2.

Fig. 3. Corrosion–wear data after reciprocation sliding tests in 3% NaCl.

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2.3. Corrosion–wear tests Corrosion–wear tests were carried out at ambient temperature (18–20 ◦ C) using two types of test: (i) reciprocation sliding against an Al2 O3 ball, and (ii) unidirectional sliding contact tests against a WC–6% Co ball. In the former case, tests were made in 3% NaCl solution, whilst in the latter, both 3% NaCl and 6% FeCl3 solutions were used. For the unidirectional tests, coated and uncoated 316L (88 mm diameter) discs were used, whilst for the reciprocation sliding tests, coated and uncoated 316L test plates, approximately 60 mm × 20 mm × 10 mm were prepared. All test surfaces were polished to a mirror finish, using 1 ␮m diamond paste, prior to coating. Uncoated test surfaces were similarly polished. Other test details are summarised in Table 2.

3. Results 3.1. Crystal structure XRD revealed that all the coatings were crystalline. All the S-phase variants were entirely single phased. This is in keeping with them having [N] within the S-phase range of 7.9 < [N] < 32 at.% N. The CrN coatings comprised CrN, with a strong {2 0 0} texture, together with traces of Cr2 N. 3.2. Electrochemical scratch tests The average corrosion current produced during electrochemical scratch testing, using an inert Al2 O3 stylus, are

shown in Fig. 2. These results demonstrated that relatively high corrosion currents developed during sliding contact between Al2 O3 and unprotected 316L. This is caused by disruption to the 316L passive layer. The tests also showed (Fig. 2) that S-phase coatings develop a lower average corrosion current, compared to uncoated 316L, and that the average corrosion current during scratching decreased with increasing [N] of the S-phase coatings. Hence, S-phase(2) coated 316L produced a lower sliding contact corrosion current than S-phase(1) coated 316L. 3.3. Reciprocation corrosion–wear tests These tests showed that S-phase coated 316L was more resistant to corrosion–wear than uncoated 316L (Fig. 3). Further, resistance to corrosion–wear increased with increasing [N] for the S-phase coatings. Examination of the corrosion–wear tracks revealed smoothly worn surfaces (Fig. 4) coupled with some evidence of superficial surface deformation. 3.4. Unidirectional corrosion–wear tests Unidirectional corrosion–wear tests using a WC–6% Co slider, yielded more complex results. Here, S-phase(3) (23 at.% N) coated 316L proved more resistant than untreated 316L. However, CrN coated 316L proved even better (Fig. 5). Surface roughness measurements (using 2-D contacting profilometry) indicated that the WC–6% Co test ball contact areas became roughened during corrosion–wear tests. The extent of this roughening was worst when the

Fig. 4. Track surfaces after reciprocation sliding corrosion–wear testing in 3% NaCl solution with Al balls: (a) S-phase(1) coating—10% N; (b) S-phase(2) coating—25% N.

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discs. This coupling was strongest for the S-phase coated test disc. Examination of the S-phase coated discs after testing showed a correlation between the wear track roughness and the worn ball surface (Fig. 8). This indicates that the coated wear surfaces became worn through a process of abrasion, in addition to the prevailing corrosion interactions. Tests conducted in 6% FeCl3 were more aggressive than those carried out in 3% NaCl. These caused failure of the coatings to take place via blistering, which could be observed within and outside the sliding contact area (Fig. 9). After forming within the corrosion–wear track, the tops of the blisters frequently became removed via a process of localised fracture.

4. Discussion Fig. 5. Corrosion–wear data after unidirectional testing against WC–6% Co balls in 3% NaCl.

WC–Co ball slid against the S-phase(3) coating than against the CrN coating (Fig. 6). This was true in both 3% NaCl and 6% FeCl3 , although the extent of roughening was greatest in the latter solution (Fig. 7). The roughening was evidently caused by corrosion attack of the Co phase of the WC–Co balls due to a galvanic coupling between the balls and the test

Corrosion–wear processes are complex because synergy takes place between the individual corrosion and wear mechanisms. Mechanical wear mechanisms are accelerated by prior corrosion mechanisms and vice versa. Corrosion–wear is an issue in bearings and valves exposed to aqueous corrosive solutions. Such environments are frequently found in food processing equipment, marine components and in bio-medical devices such as human joint replacements

Fig. 6. WC–6% Co ball surfaces: (a) after sliding against CrN coating; (b) surface roughness profile of (a); (c) after sliding against S-phase(3) coating; (d) surface roughness profile of (c). All tests made in 3% NaCl solution.

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Fig. 7. WC–6% Co ball surface after sliding against S-phase(3) coating in 6% FeCl3 : (a) light-optical micrograph; (b) 2-D surface roughness profile.

where oil-based lubricants cannot be used. Despite the wide significance of corrosion–wear, comparatively little research has been undertaken on this subject, in comparison to the quantity of investigations on mechanical wear processes and pure corrosion. Certainly, there is a dearth of research on how corrosion–wear processes can be mitigated through surface engineering. The work reported here together with other recent research at Leeds shows that the corrosion–wear of surface engineered metals and alloys is complex. Overall, there are three important corrosion–wear mechanisms that affect 316L stainless steels when they are protected by hard cathodic surface coatings like S-phase and CrN: • Type I: The removal of the coating passive film during sliding contact and its subsequent regeneration. • Type II: Galvanic attack of the substrate—leading to blistering and fracture followed by the removal of coating fragments during sliding contact. • Type III: Galvanic attack of the counterface material which causes it to roughen—this leads to mechanical damage (abrasion) of the coating during subsequent sliding contact. Type I corrosion–wear is schematically depicted in Fig. 10. This was the most probable cause of the smooth

Fig. 8. Corrosion–wear damage of WC–6% Co ball surface after unidirectional testing in 3% NaCl sliding against S-phase(3) coated 316L: (a) high magnification view of ball surface, SEM; (b) surface roughness profiles of ball and corresponding S-phase(3) coated 316L after testing.

wear processes that were rate controlling during the reciprocation sliding contact tests (Fig. 4). The electrochemical scratch test results (Fig. 2) showed that the corrosion current developed was higher for uncoated 316L than for both S-phase(1) and S-phase(2) coated 316L. Further, the corrosion current was inversely proportional to the [N] of the S-phase coating. Hence, Type I corrosion–wear is mitigated by the selection of a high [N] S-phase coating (Fig. 3). The unidirectional corrosion–wear tests illustrated that two other corrosion–wear mechanisms are important: Types II and III. These are schematically depicted in Figs. 11 and 12. Type II corrosion–wear depends on the intensity of galvanic coupling (corrosion current density) between the coating and the substrate and on the Cl− ion concentration of a given aqueous solution. Unidirectional tests in 6% FeCl3 caused blistering of the S-phase(3) (23 at.% N) coating to take place. Those blisters that were formed within the test track became fractured (Fig. 9) in the manner depicted in Fig. 11. This is a highly destructive form of corrosion–wear

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Fig. 9. Types II and III corrosion–wear after 2 h unidirectional corrosion–wear testing in 6% FeCl3 solution: (a) ruptured blister in S-phase(3) coating outside corrosion–wear test track—coating locally removed; (b) inside S-phase(3) test track—damaged blister and deep scoring due to abrasion; (c) as (b) but a smaller blister is shown with partly removed coating; (d)–(f) CrN blisters within test track with locally removed coating and light abrasion marks. Small arrows indicate direction of WC–Co ball slider travel.

which must be avoided in practice. There was a lower incidence of blistering of the CrN coated 316L, although it did take place. Presumably, in this case, there was a less intense galvanic coupling between the CrN coating and its substrate. Type III corrosion–wear (Fig. 12) also took place during the unidirectional sliding contact tests. The WC–6% Co slider (ball) material was electrochemically active—the Co phase being preferentially attacked. This subsequently caused the WC grains to drop out resulting in a roughening of the ball surface (Fig. 8). Once the WC–Co ball surfaces become roughened, they act rather like cutting tools and score the coated discs through abrasion (Fig. 12). Type III corrosion–wear took place in both 3% NaCl and 6% FeCl3 solutions (Figs. 6–9) but was worst in the latter solution. Type III corrosion–wear can be best mitigated by avoiding galvanic coupling between the sliding surfaces. For the reciprocation sliding tests, the Al2 O3 ball counter-

face was electrochemically inactive. Consequently, no adverse galvanic corrosion coupling took place and Type III corrosion–wear was never observed. The overall corrosion–wear rates in 3% NaCl during the unidirectional tests are shown in Fig. 5. This showed that CrN provided better protection than the S-phase(3) coating (23 at.% N). It appears that the corrosion attack of the WC–6% Co ball was less intense in the former case, the ball surface being smoother (Fig. 6). This indicates there was a less intense galvanic coupling between the WC–Co and the CrN coating, than was the case between the S-phase(3) coating and WC–Co. In the former case, this mitigated abrasion resulting from the Type III corrosion–wear mechanism. Also, the higher hardness of the CrN (Table 1) provided greater protection from abrasion. From a design point of view, it can be appreciated that the selection of appropriate hard coatings for corrosion–wear

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Fig. 10. Type I corrosion–wear of a hard coated metallic alloy like S-phase coated 316L.

applications is not simple. This includes, for example, the sliding contact surfaces of food filling valves, the bearing surfaces of articulating orthopaedic implant as well as the working surfaces of various marine bearings and transmission couplings. As a first priority, mitigation of galvanic coupling between the coated surface and the counterface surface is essential—this avoids Type III corrosion–wear. Secondly, the minimisation of galvanic coupling between the coating and its substrate is highly desirable. This becomes increasingly important as the chloride ion concentration of the aqueous corrosion medium increases. In this way, Type II corrosion–wear can be minimised. Lastly, a coating

Fig. 12. Type III corrosion–wear of hard coated 316L. M+ denotes generic metal ions.

with a high passive film strength and toughness (resistance to disruption) and a low corrosion activity will minimise Type I corrosion–wear. However, these are ideals that may not be always feasible for a given corrosion environment, since a high resistance to Type I corrosion–wear may mean a greater susceptibility to Type II corrosion–wear.

5. Conclusions The following conclusions are based upon sliding contact corrosion–wear tests carried out on a series of CrN and S-phase coated 316L test-pieces, immersed in various aqueous media:

Fig. 11. Type II corrosion–wear of a hard coated metallic alloy. Pitting/blistering of the hard coating culminates in mechanical fragmentation and removal (sequence (a)–(d)).

1. Three types of corrosion–wear process can affect 316L austenitic stainless steels protected with thin (5–10 ␮m) hard cathodic coatings like S-phase and CrN: • Type I: The removal of the coating passive film during sliding contact and its subsequent regeneration. • Type II: Galvanic attack of the substrate—leading to blistering and then removal of the coating during sliding contact. • Type III: Galvanic attack of the counterface material causing it to roughen—this leads to abrasion of the coating during subsequent sliding contact. 2. Type I corrosion–wear of S-phase coatings is reduced as the [N] concentration of the coating is increased. 3. Types II and III corrosion–wear processes of S-phase and CrN coatings increases with increasing chloride ion concentration in the corrosive fluid. 4. Selection of an electrochemically inert counterface material, like Al2 O3 , mitigates Type III corrosion–wear.

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Acknowledgements The authors wish to thank the Engineering and Physical Sciences Research council for the provision of a research grant (GR/L92341) and a Ph.D. studentship that made this research possible when Dr. Aldrich-Smith was at The University of Leeds. They are also pleased to acknowledge the assistance of Dr. Karl Dahm (University of Leeds) who contributed to aspects of the experimentation. References [1] R.B. Waterhouse, Fretting Corrosion, Pergamon Press, London, 1975, pp. 186–189. [2] R. Oltra, in: A.A. Sagues, E.I. Meletis (Eds.), Wear–Corrosion Interactions in Liquid Media, TMS, Warrendale, PA, 1991, pp. 3–18.

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[3] P.A. Dearnley, A. Namvar, G.G.A. Hibberd, T. Bell, Some observations on plasma nitriding austenitic stainless steels, in: Plasma Surface Engineering, vol. 1, DGM, Oberursel, 1989, pp. 219–226. [4] Z.L. Zhang, T. Bell, Surf. Eng. 1 (1985) 114–130. [5] K. Ichii, K. Fujimura, T. Takase, Structure of the ion nitrided layer of 18-8 stainless steel, Technological Report No. 27, Kansai University, March 1986, pp. 135–144. [6] K.L. Dahm, P.A. Dearnley, S-phase coatings by unbalanced magnetron sputtering, Surf. Eng. 12 (1) (1996) 61–67. [7] K.L. Dahm, P.A. Dearnley, On the nature, properties and wear response of S-phase (nitrogen-alloyed stainless steel) coatings on AISI 316L, Proc. Inst. Mech. Eng. 214 (L4) (2000) 181–198. [8] O. Vingsbo, S. Hogmark, B. Jonsson, A. Ingemarsson, in: P.J. Blau, B.R. Lawn (Eds.), Microindentation Techniques in Materials Science and Engineering, STP 889, ASTM, Philadelphia, 1986, p. 257. [9] G. Aldrich-Smith, Ph.D. Thesis, University of Leeds, 2000.