SmJhce and Coatings Tectmology, 54/55 (1992) 557--562 557

Nanoindentation and microindentation studies of hard carbon on 304 stainless steel

Teresa W r i g h t and T r e v o r F. Page Materials Division, Department of Mechanical, Materials and Manufacturing Engineering, Herschel Building, The University, Newcastle upon Tyne, NEI 7RU (UK)

Abstract

Microhardness and nanoindentation techniques have been used to probe the surface and near surface mechanical properties of hard carbon (a:C-H) coatings (0.1-t I-tm thick) deposited by an ion beam method on 304 stainless steel substrates. At low loads, nnnoindentation experiments clearly delineate the differing elastic and plastic properties of the coatings as a funetlon of thiekne3s and allow the fine-scale indentation-induced deformation and damage mechanisms to be assessed. By contrast, mieroindentation experiments have been used to show the cracking and deformation behaviour at larger contact loads where the behaviour of the substrate is important. The experimental techniques will be described, the results compared and the implications presented and discussed.

1. Introduction

If the benefits of thin hard coatings are to be fully exploited in improving the surface properties of engineer- ing components, then we need to understand more fully the detailed mechanisms by which surface properties are changed and enhanced. This understanding can be sup- plied either through modelling the response of coated systems [-11 or, experimentally, by determining the changes conferred on the system by the coating. In the latter case, it is critical that property measurements are made with the coating in situ on the test sample (i.e. in its processed state of microstructure, residual stress and adhesion) and at scales whereby the changes can be clearly recognised [2]. Recent modelling work [1] has suggested not only that the coating can modify the load eventually transmitted to the substrate (by part of the load being supported by "membrane stresses" developed as the coating is stretched) but also that the mechanical properties of the substrate are critical in determining overall responses. Thus the mechanisms by which hard carbon enhances the properties of silicon may not be the same as those by which more ductile substrates are improved.

This paper describes studies of the surface deformation of thin hard carbon films on a ductile stainless steel substrate. The principle investigative technique has been nanoindentation (using loads of 0-10mN) enabling changes to the elastic, plastic and fracture properties of the film-substrate combination to be characterised at contact displacements up to the order of the coating thickness (0.1-1 gin) where the properties of the coating

itself should dominate the contact response. Comparison has also been made with the microindentation response (~0.25-5N loads) so that changes as a function of increasing contact load--and increasing substrate importance--can be investigated.

2. Experimental details

Samples comprised hard carbon films of various thick- ness (0.1-1 gm) deposited on 304 stainless steel by the ion-beam deposition method of Franks et al. [3]. Prior to deposition, the substrates were polished, thoroughly ultrasonically cleaned and, finally, argon ion bombarded prior to deposition in the same chamber using acetylene as the base process gas.

Ultra low load indentation (nanoindentation) experi- ments were carried out using a Nano Indenter TM I[ mechanical properties microprobe (Nano Instruments Inc., Knoxville, TN, USA). Trigonal (Berkovich) inden- tations were made (under both load and displacement control) in regular arrays, with 25 ~tm between impres- sions. Constant load "hold" segments were always used (at near-complete unload) to allow any thermal drift of the instrument to be measured and corrected. The same positions were located using a software-driven stage of CamScan S4-80DV scanning electron microscopy (SEM) where the indentations were imaged in secondary electron mode using small probe sizes and mid-range kV settings to optimise image resolution. SEM imaging of nanoindentations is a challenging and time consuming process, but is needed if the load-displacement data from

0257-8972/92/$5.00 �9 1992 -Elsevier Sequoia. All rights reserved

558 T. Wright. T F. Page / Nanoimlentation aml micro�9149 studies

a nanoindentation cycle are to be interpreted to reveal detailed deformation mechanisms [-4, 5].

Microindentation testing was carried out at room temperature using a Shimadzu model "M" microhard- ness tester with a Vickers profile indenter and loads in the range 25-500 gf. Again, the appearance of the inden- tations was investigated using both reflected light micros- copy (RLM) and SEM techniques.

In all cases, the indentation response of the substrate alone was also characterised so that the changes con- ferred by the coatings could be determined.

3. N a n o i n d e n t a t i o n r e s p o n s e

Figure l(a) shows the load-displacement responses of the 304 stainless steel substrate alone for three peak

loads of 1, 5, and 10 naN. The response is typical of a ductile material, i.e. a plasticity dominated, parabolic loading curve with little elastic recovery of tlle depth displacement on unloading. Plastic deformation (i,e. non-coincidence of the loading and unloading curves) is evident even at loads as low as 1 mN (inset).

Figures l(b)- l(d) show similar plots for samples coated with film thicknesses of 0.1, 0.5, and 1 I-tm respectively and taken to the same peak loads. Compared to the substrate alone, the following changes are observed in all cases:

(1) the displacement at each peak load is reduced, the effect increasing with increasing coating thickness (i.e. surface displacement resistance has increased);

(2) the load-displacement curves exhibit increased amounts of elastic recovery on unloading, establishing that a higher proportion of the deformation is accommo-

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Fig. I. Nanoindentation load-displacement plots for (a) 304 stainless steel substrate only; (b) 0.1 Ix�9 coated steel; (c) 0.5 gm coated steel; (d) I lira coated steel. In all cases, the same three peak loads (I raN, 5 mN and 10 raN) have been used. The response curves for the 5 mN and 10raN loads are shown superposed while that for I mN is shown inset separately. Essentially, the hard carbon coatings increase the proportion of elastic response to surface contact, the effect increasing with increasing coating thickness (see text)�9

T. Wright, 7~ F. Page / Nanoindentation and microimtentation studies 5 5 9

dated elastically. Again this effect increases with coating thickness such that, while barely detectable for the 0.1 l-tin coating, the unloading curve for the 0.5 gm and 1 Ixm cases very clearly recover to smaller displacement on unloading. Indeed, in the 1 mN peak load cases for the 1 gm coating, the unloading curve retraces the loading curve exactly to show a now completely elastic response. Further, as the unloading curves become less steep (showing a larger proportion of elastic recovery of the indenter displacement/, very significant levels of elastic recovery occur during the final stages of unloading and are perhaps driven by the drum-skin effect of the mem- brane (and residual) stresses in the coating;

(3) in many cases the loading curves show steps (e.g. "S" in fig. l(e)) which, in our previous work, have been associated with the occurrence of both through thickness and interfacial cracks during flexure of the coating into the substrate.

The high resolution SEM micrographs (Fig. 2) clearly show the trigonal profile of the indenter, establish that some permanent triangular impressions remain after unloading and show the various types of crack geometry found to occur even at these low contact loads. Interes- tingly, clear differences are observable between the detailed impression left within the cracked portions of the coating and the remainder which appears less dis- tinct. This is believed to be due to elastic recovery driven by the stresses within the coating. Within the cracked regions, these stresses can have little or no effect and a larger proportion of the total deformation remains on unloading. The rippled appearance of the substrate (owing to etching of individual grains and grain bound- aries) is still visible through the coating and demonstrates how substrate topography can be replicated on the coating surface.

Within the indentations, two types of cracking are observed. Most indentations displayed a Y-shaped crack where the coating has spiit across the edges of the indenter. Additionally, many indentations display further cracks as arcs between the arms of the "Y" cracks, often leading to small areas of the coating being separated from the rest. Interestingly it is always these areas which appear to have "sunk" into the plastically depressed substrate suggesting that these latter cracks either occur as the substrate yields and plastically deforms under the coating (e.g. the behaviour is analogous to the cracking of ice on mud, or '~creme brfil6e") or occur during unloading as the membrane stresses within the coating drive recovery and attempt to pull the coating up off the deformed substrate.

In all these cases it must be emphasised that the indenter does not cut through the coatings but bends and flexes them into the substrate which, in turn, deforms under the influence of the transmitted stresses i-5, 6]. Thus the nanoindentation properties at a given displace-

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ment do not correspond to properties at a given depth into the coated surface: rather to properties with the surface flexed to this displacement:

The nanoindenter can also be used to investigate the creep behaviour of the films. Such experiments are seen as important in helping to determine the extent to which these coatings (like many other amorphous hard carbon materials) exhibit properties characteristic of stiff polymers (e.g. visco-elasticity) [3, 5, 7]. Narioindentation creep experiments were carried out by .indenting the

560 T l, Vright, T. F. Page / Nanoindentation and microindentation studies

thickest (1 gm) coating under load control to a displace- ment of 250 nm (to minimise any substrate contribution to the behaviour) and then holding the indenter at contact load for 1000 seconds while monitoring the displacement. Near complete unloading, another "hold segment" was performed to allow the thermal drift of the system to be calculated. Figure 3 shows a typical time-displacement characteristic for this coating. The calculated creep displacement rate is 0.044 nm s- t while the thermal drift rate of the system was averaged at .~" 0.01 nm s - t the highest value being 0.016 nm s- t . This suggests that time-dependent deformation of these coat- ings will be important both in accommodating surface deformation and its subsequent recovery. Thus this effect might be expected to alter the appearance of indentations in the time between testing and SEM observation. That details of cracks etc. remain distinct suggests that, long after testing, time-dependent recovery has not occurred in these regions and thus that creep is controlled by the stresses within the undamaged regions of the coating. Further, if through-thickness cracks are formed in the coating during the initial loading phase of the indenta- tion creep experiment (as seems likely given the fracture observations reported earlier), then the region of the coating directly under the indenter cannot behave like a continuous stressed membrane and, again, this suggests that the observed creep must be controlled (at least in part) by either the stretched region of the coating outside '~ the contact zone or perhaps even the substrate. Experi- ments are in progress to clarify these possibilities. ~5

4. Microindentation response

Figure 4 shows a plot of Vickers hardness number (VHN) os load for the substrate alone and with the three different coating thicknesses. Even at displacements which are now large compared to the coating thickness, the coatings enhance the hardness of the substrate. As expected, this effect is greatest for the thickest coatings and the lowest loads. Also, the coated samples display a pronounced indentation size effect (ISE) (i.e. the effec- tive hardness increases as the contact load is decreased) as would be expected from the increasing contribution of elastic response to indentation which will be most marked at lower contact loads (as confirmed by the nanoindentation studies).

The incidence of coating fracture is now much more pronounced. Networks of cracks appear in the coating and are presumed to result from the coating being unable to stretch to conform either to the increased surface area demanded by the sloping faces of the pyramidal indentation (about 8% strain compared to the original flat surface) or to the displacements created by plastic deformation of the substrate (e.g. upthrust of

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plastically displaced material around the contact area), From the SEM images (Fig. 5), it is seen that the coating has not only cracked but that the cracks have been pulled open by plastic flow of the underlying substrate. For the thicker coatings there are also regions where areas of the coating have been completely removed (Fig. 6).

It is known that TiN coatings on steels, have regular arrays of nested parallel cracks on the sides of the indented region [8] and these are also seen here. These arise during loading as the indenter bends and stretches the coating into the substrate, creating tensile stresses on the coating around the edges of the indentation (see Fig. 7). The cracks then reveal successive positions of

T. Wright, T. F. Page / Nanoindentation and microindentation studies 561

Fig. 5. SEM micrograph of a Vickers microhardness indentation in a 0.5 ~tm thick hard carbon coated steel, showing cracking (r and parting (p) of tile coating.

Fig. 6. SEM micrograph of a Vickcrs microhardness indentation in a 0.5 gm thick coating showing a region where the coating has been removed from the substrate.

] Region of Coating

S u b ~ t e ~ Through-thickness cracks

- - - ' ~ S ubstrate deformation

Fig. "7. Schematic diagram of the indentation-induced cracks which form in a brittle coating-ductile substrate system. Through-thickness fracture is largely produced by tension and, in the microhardness case, can form a series of cracks nested within the indentation outline.

the indentation edges and can be suppressed if very large residual stresses are present [-6]. Similar behaviour occurs for a: C-H on 304 stainless steel.

That the cracks observed here do not always form straight lines could be owing to local variations in the fracture strength of the coating, the adhesion between the coating and the substrate being variable (and depen- dent upon the microstructure of the substrate) or from the large degree of surface roughness of the sample resulting in a variable friction between the indenter and the surface.

Also of interest are the regions along the internal edges of the indentation where the coating appears to be cracked but with the crack edges abutting and deformed to point upwards. This is assumed to occur on unloading as elastic recovery within the coating (or substrate) recompresses the stretched]cracked coating along the indentation apices, resulting in local upward buckling and deformation of the crack edges. The impli- cation is that the coating itself displays some ductility but not enough to prevent fracture as the coating is stretched during indentation. This contrasts with the behaviour of thicker, residually stressed TiN and TiC coatings where sufficient ductility occurs to prevent any obvious fracture during indentation at room temperature and low loads [6, 8].

5. Conclusions

(1) Nanoindentation tests clearly show that even very thin a: C-H coatings on a ductile substrate improve the surface deformation resistance.

(2) At least in part, this is caused by the coating causing an increase in the contribution of elastic defor- mation to the contact response. The effect is more pronounced at higher coating thicknesses and comp]ete elastic recovery of contact displacements can be observed at very low loads.

(3) Even at contact loads of less than 10 raN, some through thickness cracking of the coating occurs, result- ing in isolated sections often appearing to be left sunk below the level of the rest. These sections usually display the sharpest detail in SEM images. The implication is that such sections have not undergone the same elastic recovery as the rest of the coating and are still attached to the plastically deformed substrate. However, whether the rest of the coating has pulled away from the substrate during elastic recovery or whether it has re-deformed, the substrate is unknown at present.

(4) Membrane and residual stresses seem to govern much of the displacement recovery.

(5) The enhanced contact resistance is still seen in the microhardness range with the increased elastic recovery

562 T. Wright, 7: F. Page / Nanoindentation and microindentation studies

leading to a large observed ISE where the hardness increases at lower contact loads.

(6) Cracking is more widespread in the micro- indenta t ion tests with nested cracks appea r ing within the indenta t ion outlines.

(7) In all cases, S E M is invaluable in revealing the deformat ion which has taken place.

(8) Nano inden ta t ion has been used to investigate creep of the coating, p resumably by a viscous mechanism within the "polymer-l ike" a: C - H film with a creep ra te of 0.04 nm s - t being obtained at cons tan t loads of 10raN.

Acknowledgments

The Materials Commission of the Uni ted K i n g d o m Science and Engineering Research Counci l (SERC) are thanked for providing the Nano Inden te r TM II used for these experiments and for the ceramics SEM facility.

T.W. would like to thank SERC and Ion Tech Ltd. for financial suppor t . Dr A. J. Whi t ehead is t h a n k e d for useful discussions.

References

1 P.M. Ramsey, H. W. Chandler and T. F. Page, Sm:f. Coat. Technol., 49 ([991) 504.

2 J. C. Knight and T. F. Page, Proc. 2nd European Colloquium on Designing Ceramic Interfaces, Petten, The Netherlands, November, 1991, Elsevier, in the press.

3 J. Franks, K. Enke and A. Richardt, Met. Mater., 6 (1990) 695. 4 T. F. Page, W. C. Oliver and C. J. McHargue, J. Mater. Res., 7

(1992) 450. 5 J. C. Knight, A. J. Whitehead and T. F. Page, J. Mater. Sci., 27

(1992) 3939. 6 T. F. Page and J. C. Knight, Sulf. Coat. Technol., 39/40 (1989) 339. 7 J. C. Knight, T. F. Page and 14. W. Chandler, S m f Coat. Technol.,

49(1991) 519. 8 J. C. Knight, T. F. Page and I. M. Hutchings, Surf. Eng., 5 (1989)

213.