Microstructural effects on the friction and wear of zirconia films in unlubricated sliding contact

Microstructural effects on the friction and wear of zirconia films in unlubricated sliding contact

Thin Solid Films 347 (1999) 220±225 Microstructural effects on the friction and wear of zirconia ®lms in unlubricated sliding contact S.C. Moulzolf a...

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Thin Solid Films 347 (1999) 220±225

Microstructural effects on the friction and wear of zirconia ®lms in unlubricated sliding contact S.C. Moulzolf a,*, R.J. Lad a, P.J. Blau b a

Laboratory for Surface Science & Technology, University of Maine, Orono, ME 04469-5764, USA Metals and Ceramics Division, Oak Ridge National Laboratory, Oak Ridge, TN 37831-6063, USA


Received 8 November 1998; received in revised form 7 January 1999; accepted 7 January 1999

Abstract Friction and wear of sapphire and steel counterfaces sliding on unlubricated thin ®lms of zirconia were studied using conventional multiple pass pin-on-disk testing and single pass testing with a `friction microprobe'. The zirconia ®lms were grown by electron-cyclotron-resonance (ECR) oxygen plasma assisted deposition on r-cut (011Å2) sapphire using different growth parameters to produce ®lms with varying degrees of order. The microstructure of the zirconia ®lms is correlated to the widely varying observed friction and wear characteristics. Anisotropic stress effects produced during scratch tests are also presented. q 1999 Published by Elsevier Science Ltd. All rights reserved. Keywords: Zirconium oxide ®lm; Friction; Wear; Microstructure; Sliding contact

1. Introduction Zirconia has been widely studied due to its desirable physical properties such as a high melting point, low thermal conductivity, high coef®cient of expansion [1] and resistance to corrosion [2,3]. Bulk zirconia exhibits three equilibrium polymorphs: monoclinic to ,11008C, tetragonal to 23708C, and cubic up to the melting point of 26808C [1]. Transformation from tetragonal to monoclinic phase during thermal cycling is accompanied by a volume expansion of 3% to 5% which can lead to failure by cracking. This large volume change can be suppressed by stabilizing zirconia in either the cubic or tetragonal phase [4] by addition of a suitable oxide dopant such as yttria, calcia or magnesia [5]. By appropriate oxide doping and processing to produce ®ne grained ceramics, metastable tetragonal zirconia polycrystals can be synthesized. This tetragonal phase improves the fracture toughness and strength of bulk zirconia through a tetragonal-monoclinic transformation in the presence of a stress ®eld [4]. The large volume of reported work on wear behavior of bulk zirconia reveals varied results depending on microstructure and composition. For example, moisture has been reported to have a deleterious effect on bulk yttriastabilized tetragonal zirconia leading to transformation to * Corresponding author. Tel.: 1 1-207-581-3043; fax: 1 1-207-5812255. E-mail address: [email protected] (S.C. Moulzolf)

monoclinic phase and microcracking at low temperatures ([6] and refs. cited therein). Fischer et al. [7] observed different wear behavior for yttria-stabilized cubic and tetragonal zirconia in different environments (nitrogen, air, water, hexadecane). Also, Lancaster et al. [8] present a summary of the qualitative effects of water on wear for many reported bulk zirconia systems. There is a relatively small amount of work reported on friction and wear testing of zirconia thin ®lms. Friction and wear studies have been performed on zirconia ®lms deposited by various techniques such as sol gel coating [9], chemical vapor deposition [10], sputter deposition [11,12] and plasma spraying [13]. Much of this work on friction and wear of zirconia coatings has been carried out in the context of rigid disk magnetic recording media. Yamashita et al. [11] found that unlubricated sputtered yttria-stabilized coatings having thicknesses between 20 and 30 nm maintained static friction coef®cients less that 0.4 for repeated contact start stop (CSS) tests using a read/write head with a force of 0.15 N. Under similar testing, Ganapathi et al. [13] observed cracks perpendicular to the ®lm surface running parallel to the sliding direction after 5000 CSS cycles using a load of 0.093 N. This cracking was attributed to compressive stress in the ®lm. Dugger et. al. [12] report that 30-nm thick sputter-deposited yttria-stabilized zirconia overcoats lubricated with per¯uoropolyether have greatly improved friction and wear characteristics in humid air compared to dry air or vacuum environments. Oxidation of metallic wear debris

0040-6090/99/$ - see front matter q 1999 Published by Elsevier Science Ltd. All rights reserved. PII: S00 40-6090(99)0004 6-2

S.C. Moulzolf et al. / Thin Solid Films 347 (1999) 220±225

Fig. 1. RHEED images from zirconia ®lms at 30 kV accelerating voltage. (a) `Random' polycrystalline zirconia; (b) `highly oriented' polycrystalline zirconia (c) `oriented/random' polycrystalline zirconia.

and coating alteration by water vapor were proposed as mechanisms for the improved characteristics. This paper presents a study of the effects of microstructure on the tribological performance of zirconia thin ®lms on sapphire substrates prepared by electron cyclotron resonance (ECR) oxygen-plasma-assisted deposition. In contrast to typical sintered bulk zirconia, the ®lms in this study possess well de®ned surface microstructure and contain no dopants or impurities. We have previously reported that highly oriented ®lms containing cubic phase zirconia can be synthesized on (011Å2) a -Al2O3 single crystal substrates (r-cut sapphire) without addition of stabilizing dopants [14]. Friction and wear properties of this type of zirconia ®lm are reported and compared with other undoped zirconia ®lms possessing varying degrees of order which were grown on r-cut sapphire. 2. Experimental The MBE thin ®lm growth system, in situ composition and structural analysis capabilities, and zirconia ®lm growth parameters are described in detail in a previous paper [14]. Brie¯y, the zirconia ®lms were grown in an ultrahigh vacuum chamber by electron beam evaporation of Zr metal in the presence of an oxygen plasma generated by a multi-pole ECR plasma source. The substrates were commercial epi-polished r-cut sapphire wafers which were etched in a hot phosphoric acid solution to remove any polishing residue and subsequently annealed at 10008C for 90 min in 5 £ 105 Torr oxygen to remove hydrocarbons and order the surface. This treatment yields a clean highlyordered stoichiometric (1 £ 1) surface as determined by X-ray photoelectron spectroscopy (XPS) and re¯ection high energy electron diffraction (RHEED). RHEED and atomic force microscopy (AFM) were utilized to characterize ®lm microstructure and topography. Tribological characterization of the zirconia ®lms was carried out at the High Temperature Materials Laboratory at Oak Ridge National Laboratory. Friction and wear characteristics were determined using a pin-on-disk (POD) apparatus with the sphere on rotating ¯at con®guration in ambient air with a relative humidity of ,80% as measured using the wet/dry bulb method. Both 440C stainless-steel and sapphire spheres of 10 mm diameter were used as counterfaces. All tests were performed with a 1 N normal force,


0.1 m/s speed, and total sliding distance of 10 m. All tests were repeated under identical conditions on separate regions of the ®lms to check for reproducibility. Frictional force was measured using a load cell arrangement and monitored on a chart recorder. Optical microscopy, pro®lometry and AFM were used to study wear characteristics of the ®lms. The friction microprobe (FMP), a specialized tribometer developed at Oak Ridge National Laboratory, was utilized to study the friction characteristics of the ®lms in the microcontact regime. Details of the FMP and applications of its use on other systems are available elsewhere [15±17]. Sapphire and 440C stainless-steel spheres were used for comparison with pin-on-disk results, but the spheres in these tests were 1 mm in diameter. Single-pass tests of stroke 1000 mm were performed in ambient air at 10 mm s 21 with a normal force of 98.1 mN. Frictional force was measured using a capacitive gap sensor and data were digitally stored at a 20 Hz sample rate. A Revetest scratch tester by CSEM was used to study ®lm failure. The scratch tester was used to generate a controlled scratch on the ®lm at a constant load using a 200 mm radius of curvature conical diamond tip at a rate of 0.2 mm/s. 3. Results and discussion 3.1. Microstructural characterization The zirconia ®lms were grown on r-cut sapphire substrates at varying temperatures and rates to control the ®lm microstructure. Films grown at temperatures below 3008C at rates from 0.05 to 0.3 nm/s had a random polycrystalline structure as shown by ring diffraction patterns in a RHEED image viewed from a 500-nm thick ®lm grown at 0.2 nm/s at room temperature (Fig. 1a). This ®lm had a partial ®ber texture as evidenced by incomplete rings in the diffraction pattern. Phase identi®cation was not possible because the lattice spacings for the different zirconia polymorphs are very similar and the diffraction pattern could not be indexed from the broad diffraction rings. However, recent X-ray diffraction data from our laboratory indicates that random polycrystalline cubic zirconia ®lms can be grown within the parameters stated above (unpublished data). Fig. 1b shows the RHEED pattern from a 100-nm thick ®lm grown at 4758C and 0.05 nm/s. The initial ®lm composition (up to ,30 nm thickness) was primarily highly

Fig. 2. AFM topography images (5 £ 5 mm) of zirconia ®lms imaged in noncontact mode. (a) `Random' polycrystalline zirconia; (b) `highly oriented' polycrystalline zirconia; (c) `oriented/random' polycrystalline zirconia.


S.C. Moulzolf et al. / Thin Solid Films 347 (1999) 220±225

Table 1 Summary of zirconia ®lm types with corresponding deposition parameters and ®lm properties Film type

Growth temp. (8C)

Deposition rate (nm/s)

Thickness (nm)

Phase composition

Grain size (nm)

Roughness (rms) (nm)

Random Highly oriented Oriented/random

,300 475 500

0.05±0.3 0.05 0.3

500 100 100

Random cubic Oriented cubic/monoclinic Oriented cubic/monoclinic 1 random component

90 116 163

6.7 6.9 14.4

oriented cubic phase zirconia. At greater ®lm thicknesses, highly oriented monoclinic phase zirconia began to nucleate and coexist with the cubic phase. This ®lm type will hereafter be labeled as `highly oriented' polycrystalline. Note that Fig. 1b is a transmission diffraction pattern from RHEED which indicates a three dimensional ®lm morphology. The ®lm growth occurs as three dimensional crystallites as con®rmed by AFM imaging (Fig. 2). The microstructural evolution of a ®lm grown at a temperature of 5008C and deposition rate of 0.3 nm/s was similar to the `highly oriented' polycrystalline ®lm with the notable addition of a random polycrystalline component evident from diffraction rings at the ®nal thickness of 100 nm (Fig. 1c). This ®lm type will be referred to as `oriented/random' polycrystalline. A summary of the ®lm types that were used in this study is provided in Table 1. AFM images of the unworn surfaces of the ®lms reveal different surface topographies for each of the three zirconia ®lm types that were studied. The AFM images in Fig. 2 were obtained using non-contact imaging with a Si tip of nominal radius of 10 nm; non-contact imaging was required since signi®cant tip wear during contact imaging resulted in image artifacts. Films of types `random', `highly oriented', and `oriented/random' polycrystalline had average grain

Fig. 3. Optical micrographs of sapphire counterface (above) and zirconia ®lm wear track (below) after POD testing. (a) `Random' polycrystalline zirconia; (b) `highly oriented' polycrystalline zirconia; (c) `oriented/ random' polycrystalline zirconia

sizes of 90 ^ 1, 116 ^ 5, and 163 ^ 5 nm, respectively, as estimated by the intercept method [18]. The `oriented/ random' polycrystalline ®lm type had an rms roughness of 14.4 nm over a 5 £ 5 mm area which was more than double the values for the `random' and `highly oriented' polycrystalline samples of 6.7 and 6.9 nm, respectively. 3.2. POD friction and wear characterization Pin-on-disk tests with sapphire and steel counterfaces on `random' polycrystalline zirconia ®lms resulted in ®ne-scale abrasive ®lm wear. The optical micrographs of both the ®lm wear track and the corresponding sapphire counterface after 10 m of sliding contact are shown in Fig. 3a. The ®lm micrograph shows ®lm debris at the edges of the wear track that was generated in the contact region. The adhered debris on the counterfaces was easily removed with isopropanol and no wear on either type of counterface was visible. The continuously monitored friction coef®cient for a sapphire counterface was essentially constant for the duration of the 10 m test as shown in Fig. 4a. Measured friction coef®cients for two separate tests performed under identical conditions ranged from 0.22 to 0.33 for sapphire counterfaces and 0.22 to 0.30 for steel counterfaces; the ranges indicate the minimum and maximum measured instantaneous friction coef®cient. The small oscillating component of the friction coef®cient trace has a period corresponding to one revolution of the disk specimen. The average wear coef®cient was determined from wear track cross sections measured with a pro®lometer at four spatially separated points for each of two tests. The average wear coef®cient for the ¯at specimen was calculated to be 3 £ 1025 mm 3 N 21 m 21 for both sapphire and steel counterfaces, which is an order of magnitude larger than that suggested by Czichos [19] for materials suitable for unlubricated tribological applications. Because the wear debris as observed in the optical micrographs is very ®ne, a proposed wear mechanism for this system is grain pullout. Under this assumption, the large ®lm wear rate coupled with negligible counterface wear would be indicative of low ®lm intergranular adhesion. POD testing of the `highly oriented' polycrystalline zirconia ®lm resulted in friction coef®cient values similar to those for the `random' polycrystalline ®lms, but little damage to either ®lm or counterface was produced. Optical micrographs for a sapphire counterface (Fig. 3b) show that

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Fig. 4. POD friction coef®cient traces. (a) Sapphire on `random' polycrystalline zirconia; (b) sapphire on `highly oriented' polycrystalline zirconia; (c) sapphire on `oriented/random' polycrystalline zirconia; (d) steel on `oriented/random' polycrystalline zirconia.

the POD tests produced very little debris and the contact regions of both the ®lm and sapphire counterface have a polished appearance. The ®lm wear was negligible as determined by pro®lometry. The slowly increasing friction coef®cient trace shown in Fig. 4b ranged from 0.22 to 0.44. In comparison, the friction coef®cient trace for a steel counterface (not shown) rose only slightly and actually began to decrease near the end of the test with values ranging between 0.24 and 0.30. The low wear rate for this ®lm type is likely due to low surface roughness and good intergranular adhesion between the highly oriented crystallites. An AFM image (Fig. 5a) near the center of the wear track on the zirconia ®lm produced by a sapphire counterface reveals that the worn area is relatively smooth with no large wear particles present. The visible scratches parallel to the sliding direction (vertical) are on the order of 0.1 mm which is why they cannot be resolved in the optical micrographs and the area appears polished. The `oriented/random' polycrystalline zirconia ®lm exhibited the least desirable friction characteristics for the POD testing conditions used. Sliding of sapphire and steel

Fig. 5. AFM images (5 £ 5 mm) inside ®lm wear tracks. (a) Sapphire on `highly oriented' polycrystalline zirconia; (b) sapphire on `oriented/ random' polycrystalline zirconia.

Fig. 6. FMP friction coef®cient traces for sapphire sliding on zirconia ®lms. (a) `Random' polycrystalline zirconia; (b) `highly oriented' polycrystalline zirconia; (c) `oriented/random' polycrystalline zirconia.

on this type of zirconia was abrasive and resulted in both ®lm and counterface damage. Low regions at the edges of ®lm wear tracks as measured by pro®lometry indicate that the ®lm initially wears, but a transfer ®lm of counterface material clearly forms in the center of the wear track. This transfer ®lm of counterface material and wear particulate is clearly shown in an AFM image near the center of a wear track (Fig. 5b) created with a sapphire counterface. The optical micrographs of the corresponding counterface and wear track are shown in Fig. 3c. For the steel counterface, scanning Auger electron spectroscopy (AES) detected measurable amounts of Fe in the wear track. The transfer ®lm likely occurs due to abrasive third body particulates from the random polycrystalline component of this ®lm. The `oriented/random' polycrystalline ®lm had a larger average grain size and larger rms roughness than the highly oriented ®lm. It appears that the grains in the `oriented/ random' polycrystalline ®lm are more susceptible to intergranular fracture than the `highly oriented' ®lm. The friction coef®cient traces for both the sapphire and steel counterface shown in Figs. 4c and d, respectively, reveal striking transitional behavior to high friction coef®cients. For the sapphire counterface, the measured friction coef®cients during the initial run-in period and prior to the transition point ranged from 0.30 to 0.52 and the ®nal friction coef®cient was 0.75 to 0.81 and appeared to be monotonically increasing at the end of the test. Friction coef®cients for steel counterfaces ranged from 0.29 to 0.53 before the transition point and 0.85 to 1.30 thereafter. A ringing noise was also audible at the onset of the transi-


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Fig. 7. Optical micrographs of `oriented/random' polycrystalline zirconia ®lm after scratch testing using an 8.8 N load. Micrographs are oriented to show relative scratch directions which are indicated by arrows. (a) Tensile type cracks, (b) longitudinal cracks and void formation.

tion which continued for the remainder of the test. The friction characteristics coupled with the observed counterface transfer ®lm indicates that the transition corresponds to a critical point at which the contact changes from primarily counterface/®lm to counterface/(counterface transfer ®lm). The large measured friction coef®cients after the transition point are consistent with tabulated ranges of values for unlubricated sliding of Fe/Fe and sapphire/sapphire pairs [20]. 3.3. FMP friction characterization The FMP was used to measure friction coef®cients on a smaller contact regime than POD tests. As a point of reference for comparison of the POD and FMP tests, the calculated Hertzian contact dimensions and maximum contact stress for the sapphire/zirconia system are approximately 30 mm and 10 GPa, whereas for the FMP tests they are approximately 6 mm and 100 GPa. Also, the POD test comprises multiple passes over the same area in contrast to the FMP test which is a single pass test meaning that accumulated debris is not involved in the wear mechanism. FMP friction coef®cient vs. sliding distance data for sapphire on the three types of zirconia ®lms is shown in Fig. 6. The `oriented/random' polycrystalline ®lm FMP curve (Fig. 6c) clearly demonstrates stick-slip behavior characterized by rapid oscillations in the friction coef®cient that occur with a fairly regular spacing of ,6 mm. Although all the FMP curves demonstrate large deviations in friction coef®cient, the average friction coef®cient was determined from two separate tests performed under identical conditions for comparison to the POD tests. The initial `spike' in the friction coef®cient for each FMP curve in Fig. 6, which is caused when the static frictional force for the counterface/®lm system is exceeded, was not included in the average friction coef®cient determination. Because the FMP testing involved only 1 mm of sliding distance as opposed to the 10 m distance in the POD tests, the only comparison that can be made between the tests is

between the average FMP friction coef®cient and the corresponding initial measured POD friction coef®cient. Although the POD and FMP tests involve different conditions, there was good correlation between the friction coef®cients de®ned above. The initial measured friction coef®cients for sapphire on zirconia POD tests were 0.26, 0.22 and 0.31 for `random', `highly oriented', and `oriented/ random' polycrystalline zirconia in comparison to average FMP friction coef®cients of 0.27, 0.18, and 0.30 for the same systems. The FMP curves were relatively independent of slider material, sapphire or steel, and possessed the same qualitative features and friction coef®cients within 10% for each ®lm; thus only the results for sapphire counterfaces are shown. 3.4. Scratch test characterization Scratch testing of the ®lms produced anisotropic cracking in the `oriented/random' ®lm. The tests were performed using a scratch tester with a conical diamond tip of 200 mm radius. Fig. 7a shows an optical micrograph of a scratch generated in an arbitrary direction on the oriented ®lm at a load of 8.8 N. A series of regularly spaced tensile type cracks are observed with an orientation of approximately 458 to the scratch direction. Cracks were ®rst observed at a load of 2.0 N and their number density increased with increasing load. Scratching in the direction parallel to the observed crack produces long cracks parallel to the scratch direction as shown in Fig. 7b. Also apparent in the ®gure are voids which were likely caused by delamination at the ®lm/ substrate interface. These voids were only observed for scratch testing in this particular direction. Somewhat similar cracking behavior was observed by Wu [21] during microÊ zirconia ®lm deposited on a scratch tests on a 250 A magnetic coating. However, Wu observed only semicircular cracks which were symmetric with respect to the scratch direction and which were attributed to the high tensile stress generated at the trailing edge of the scratch tip. Our scratch tests demonstrate that the cracks propagate along a preferred direction in the `oriented/random' ®lm. Because a fully random ®lm should be isotropic with respect to crack propagation, the singular crack direction suggests that the cracking primarily occurs in the `oriented' component of the ®lm. No cracking was observed in the `random' or `highly oriented' polycrystalline ®lms for loads up to 8.8 N. 4. Conclusions Unlubricated sliding on the zirconia ®lms used in this study showed a critical correlation between friction and wear characteristics and ®lm microstructure. Pin-on-disk testing revealed that `random' polycrystalline ®lms had the highest ®lm wear and friction coef®cients , 0.35. Similar friction coef®cients coupled with negligible wear on either the counterface or ®lm resulted from the `highly oriented' polycrystalline zirconia microstructure. The

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presence of a `random' polycrystalline component intermixed with a `highly oriented' microstructure resulted in friction coef®cients ,1 due to a transfer ®lm of counterface material. Macro-scale POD and micro-scale FMP tests indicate that frictional phenomena scale consistently for the materials systems studied. The `oriented/random' ®lm microstructure exhibited anisotropic cracking during scratch testing which is attributed to fracture primarily in the `oriented' component of the ®lm. Acknowledgements This work was performed under support from the U.S. Department of Energy: Grant No. DE-FG02-90ER45417; and from the Oak Ridge National Laboratory, High Temperature Materials Laboratory User Program, under contract DE-AC05-96OR22464 with Lockheed Martin Energy Research Corporation. S.C.M. is grateful for fellowship support from the DOE-EPSCoR Traineeship Program through the Maine Science and Technology Foundation. References [1] M. Bauccio (Ed.), Engineering Materials Reference Book 2, ASM International, 1994. [2] T. Yamashita, G.L. Chen, J. Shir, T. Chen, IEEE Trans. Magnetics 24 (1988) 2629. [3] I. Zaplatynsky, Thin Solid Films 95 (1982) 275.


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