Chemical and tribological studies of MoS2 films on SiC substrates

652

Surface and Coatings Technology, 43/44 (1990) 652—662

CHEMICAL AND TRIBOLOGICAL STUDIES OF MoS2 FILMS ON SiC SUBSTRATES STEPHEN V. DIDZIULIS and PAUL D. FLEISCHAUER Chemistry and Physics Laboratory, The Aerospace Corporation, P.O. Box 92957, Los Angeles, CA 90009-2957 (U.S.A.)

BONNIE L. SORIANO and MICHAEL N. GARDOS Hughes Aircraft Company, P.O. Box 902, El Segundo, CA 90245 (U.S.A.)

Abstract The lubrication of ceramic surfaces is vital to permit the use of ceramics in many types of moving mechanical assembly. This paper reports the results of pin-on-flat friction and wear studies conducted in vacuum on lubricating MoS2 films that were r.f. sputter deposited onto hot-pressed SiC coupons pretreated to alter available surface bonding sites. Pretreatment consisted of chemical etching with HF and HNO3, which removed surface active sites for MoS2 adhesion. Films grown on etched surfaces have shorter wear lives and lower initial friction than similar films grown on non-etched surfaces. The surface chemical interactions of SiC with molybdenum, sulfur and oxygen were also briefly examined using X-ray photoelectron spectroscopy. SiC was observed to form strong bonds with all three species; molybdenum had an interfacial reaction, while sulfur and oxygen competed for strong surface bonding sites.

1. Introduction The inclusion of wear-resistant, monolithic ceramic or ceramic-coated components in many types of moving mechanical assembly will require effective lubrication of the components. For space-based applications, solid lubrication is highly desirable to limit lubricant outgassing, which could result in contamination of sensitive optical components and/or mechanical failure caused by lubricant starvation. This paper presents results of studies on the growth and lubricating properties of thin molybdenum disulfide (MoS2) films on hot-pressed silicon carbide (HP SiC) substrates. The films were r.f. sputter deposited onto HP SiC coupons prepared in a variety of ways to alter available surface bonding sites. The paper also explores surface chemical interactions between SiC and molybdenum, sulfur and oxygen adatoms. 0257.8972/90/$3.50

© Elsevier Sequoia/Printed in The Netherlands

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The lubricating properties of MoS2 derive from its anisotropic, layered crystal structure. In this structure, interatomic forces between successive MoS2 sandwich layers are weak, resulting in low shear strength. The crystal structure gives the material very inert (0001) basal planes with saturated bonds and reactive edge planes with dangling bonds. Lubrication in an MoS2 film is believed to occur via intercrystalline slip along basal plane faces of adjacent crystallites [1]. Two parameters related to the crystal structure and hypothesized to be of particular importance in solid lubrication with MoS2 are the adhesion of the film to the substrate surface and the crystallographic orientation of the platelets in the film. Film—substrate adhesion is believed to occur primarily through MoS2 edge plane sites and should strongly influence the wear life of the film: the stronger the adhesion, the longer the wear life [2]. MoS2 film orientation should affect the measured coefficient of friction (COF). Low friction could presumably be obtained in films that have crystallites oriented with basal planes parallel to the plane of sliding. Initial friction measurements would probably reflect the effects of film orientation because most MoS2 films are quickly deformed to produce a surface layer of lubricating, basal-oriented MoS2 [3]. Inherent in this model of MoS2 lubrication is the notion that films having more crystallites with basal planes oriented parallel to the direction of sliding will have a lower initial friction, but will be less adherent than perpendicular films, resulting in shorter wear lives. Previous work on the MoS2/SiC system [4] has shown MoS2 film growth and morphology to be markedly dependent on SiC substrate pretreatment and on MoS2 deposition conditions. Figure 1 exhibits the relationship between the Si 2p and C is X-ray photoelectron spectroscopy (XPS) features of an HP SiC substrate, and the (002) X-ray diffraction (XRD) peak of an r.f.-sputtered MoS2 film deposited at 220 °C.This relationship is shown as a function (a~

(b)

(c)

~

~HF

ETCH CH3OH RINSE

—

106

104

102 100 BINDING ENERGY (eV)

90-’

96

I 288

I

286 284 282 BINDING ENERGY e~

_________________________________

280

10

11

12

13

14

15

16

17

2

Fig. 1. Comparison of the effects of various chemical pretreatments of HP SiC on (a) the Si 2p XPS peaks, (b) the C is XPS peaks and (c) the MoS2 (002) XRD peak of 200 non high temperature (HT) films. The chemical treatments are listed in the figure.

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of substrate treatment prior to film deposition. The XPS features indicate that the greatest effect of the chemical etches and vacuum anneal is the removal of carbon-containing contaminants. As shown in Fig. 1(c), the MoS2 (002) XRD peak grows with increasing surface cleanliness, indicating that films deposited at high temperature grow with increasing parallel basal orientation as carbon-based contaminants are removed. These results support an active site model [5] for perpendicular MoS2 film growth, where reactive surface species (carbon contaminants on SiC) act as nucleation sites for edge plane growth. It should be noted that films grown at 60 °Cexhibited no (002) XRD reflections and, hence, no dependence on surface treatment. The work presented in this paper focuses on two aspects of the MoS2/ SiC system. Firstly, results from oscillatory friction and wear tests conducted with a vacuum pin-on-flat tribotester are correlated with the SiC substrate pretreatment and Mo52 film deposition conditions. Substrate pretreatments include methanol rinses, acid etches and the deposition of a molybdenum metal interlayer. Secondly, the chemistry of SiC surfaces is explored. Included in this work are the reactions between SiC and sulfur, oxygen and molybdenum. These reactions have been monitored using XPS to determine how the SiC surface could be altered to produce higher performance lubricating films.

2. Experimental details Substrates used for the friction and wear tests were fully hot isostatically pressed ct-SiC containing 0.5% Al2 03 as a sintering agent, produced by ESK (F.R.G.). Samples were polished with successively finer diamond paste down to 0.25 j.tm grit prior to chemical treatments. Samples were chemically treated in a dry nitrogen glove bag by being rinsed with methanol (designated MeOH rinse) or etched for 1—2 s in a buffered HF solution, and then etched for 1—2 s in concentrated nitric acid. Each etch was followed by a methanol rinse (designated acid etch) [4]. Surface roughness was not measured, although scanning electron micrographs [4] show no detectable difference on the micrometer scale after acid etching. Thin molybdenum layers (nominally 10 nm) were evaporated at approximately 2 nm s’ onto some of the acid-etched SiC coupons in an electron-beam evaporator. The evaporator was contained in a diffusion-pumped bell jar with a base pressure of 2.66 x 10~Pa (2 x 10~Torr). After pretreatment, samples were stored in a vacuum (approximately 100 Pa) desiccator for less than 1 h until transfer to the r.f. sputter deposition chamber. Thin films of MoS2 were r.f. sputter deposited in a diffusion-pumped system, described in detail elsewhere [1, 6]. Films (nominal thickness 200 nm) were deposited with the substrate maintained at either the ambient deposition temperature of 60 °C (AT films) or at an elevated temperature of 220 °C (HT films). In HT film deposition, the substrate was heated for approximately 16 h in the vacuum system prior to deposition. Similar

655

films were also prepared on methanol-rinsed 17-4 PH stainless steel substrates. Single-crystal ct-SiC samples obtained from Atomergic Chemical Corporation, Plainview, New York, were used to study oxygen and sulfur chemistry. The basal planes of the SiC samples were polished with diamond paste down to 0.25 I.Lm grit. Each sample was then sputter annealed (at accelerating potentials of 1 kV, 500 V and 300 V; 550 °C) in an ion-pumped, ultrahigh vacuum (UHV) system described elsewhere [4], until no oxygen or other impurities were detected by XPS. The clean sample was then exposed to oxygen in a chamber backfilled with Matheson research purity 02. Sulfur exposures were performed in the preparation chamber with an electrochemical sulfur doser modeled after that described in ref. 7. XPS was performed in the previously mentioned UHV system with a Surface Science Instruments SSX-100, small-spot XPS system. The parameters employed for this study were a spot size of 300 ~.tmand a pass energy of 50 eV. All binding energies were referenced to the spectrometer Fermi level, with the gold 4f712 line at 84.0 eV and full width at half-maximum (FWHM) of 0.9 eV. Sputter profiling was performed with a differentially-pumped Leybold— Haraeus IQE 12/38 argon ion gun operated at 2 kV in a 2 mm x 2 mm raster mode. A pin-on-oscillating-flat friction and wear tester was used to study the tribological properties of the films. The tester, shown schematically in Fig. 2(a), was located in a turbomolecular-pumped; Cambridge Stereoscan 250 MK3 scanning electron microscope (SEM), which has been described elsewhere [8]. The operating vacuum of the system was in the 10-s Pa (10~~ Torr) regime. The pin was a 90° single crystal of A1203, which was checked for wear after each set of tests on a sample. The normal load applied to the pin was 60—65 gf, which corresponded to a 1.75 GPa maximum Hertzian stress. The were oscillated at 1 The mm s~ 15 length cycles, was and approxithen at 1forsamples the remainder of the test. pin for track 4mm s~ AFERT URE VOLTAGE PRIMARY

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SECONDARY ELECTRONS

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2~ 20 75100125 150 175 200 225 250 275 300 CYCLES

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Fig. 2. (a) Schematic diagram of the SEM tribotester. (b) Trace of the COF vs. the number of wear cycles for a 200 non AT film deposited on a methanol-rinsed SiC substrate.

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mately 3 mm. Two tests were run on each sample at room temperature. The friction force, applied load and sample temperature were measured continuously with computerized data logging. The average coefficient of friction per cycle was calculated from the real-time data. The films were judged to have failed when the COF became greater than 0.1 in the higher speed mode. While the wear tests were in progress, real-time images of the developing wear scars were recorded on videotape with the SEM using a 20 kV standard beam with a 3 mA gun current.

3. Results and analysis In Fig. 2(b), a typical trace of the COF vs. the number of oscillating cycles is shown for an AT film deposited on a methanol-rinsed HP SiC substrate. Typically, the low-speed initial breakaway COF for all films was higher than the high-speed, steady state friction (0.1 vs. 0.05 for the AT film in Fig. 2). In general, failure was easy to determine as the film delaminated, which caused a sharp increase in the measured friction and was evident in the SEM video. In Table 1, the average room temperature (RT) wear lives are presented for the MoS2/SiC and MoS2/steel films. All films deposited on SiC had significantly shorter wear lives than films deposited on steel. The latter did not fail before the tests were terminated at 1000 cycles. AT films deposited on the MeOH-rinsed and molybdenum interlayer samples exhibited longer wear lives than HT films deposited on similarly prepared substrates. Finally, the different sample preparations had a significant effect on wear lives for both AT and HT films. Films deposited on substrates rinsed in methanol exhibited TABLE 1 Room temperature wear life and friction data for the 200 non MoS2 films on HP SiC and 17-4 PH steel substrates Substrate preparation

HP SiC MeOH rinse MeOH rinse Acid etch Acid etch Mo layer Mo layer 17-4 PH steel Ground Polished

Film type

AT

Wear cycles

COF Initial

Minimum

HT AT HT AT HT

456 318 147 166 318 227

0.10 0.10 0.08 0.06 0.12 0.11

0.05 0.03 0.04 0.05 0.02 0.05

AT HT

i000~ 1000~

0.08 0.08

0.02 0.03

~Tests terminated after 1000 cycles.

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wear lives more than a factor of two greater than films deposited on the acid-etched SiC. The molybdenum interlayer films exhibited average wear lives somewhere between those of the other two films, although the individual tests varied widely for the Mo/SiC samples. Also summarized in Table 1 are the RT COF results for the first test on each sample. The COF results of only the first test are given to eliminate the effects of MoS2 film transfer to the pin, which would quite likely alter subsequent initial friction measurements. These COF results are less clear cut than the wear data. The initial COF for all films was greater than the steady state value, indicating that a run-in period is required for optimum film performance. The films deposited on steel seemed to have, in general, both lower initial friction values and lower steady state friction values than similar films deposited on the harder SiC. In comparing the various MoS2/SiC films, the most significant finding was the lower initial friction measurements on the acid-etched substrates. Both the AT and HT films had a lower initial COF than similar films on the methanol-rinsed substrates. The HT acidetched film had the lowest initial COF (0.06). The minimum steady state friction exhibited no clear trend with film type or sample treatments and ranged from 0.02 to 0.05. The friction and wear tests are consistent with the active site model discussed earlier. The most important comparison should be drawn between the methanol-rinsed and acid-etched SiC. As shown in Fig. 1, acid-etched SiC has significantly less surface carbon contamination, and hence fewer active sites, for edge-bonded MoS2. This weaker adhesion produced shorter wear live for AT and HT films deposited on the acid-etched SiC. In addition, HT films generally had shorter wear lives than AT films deposited on similarly prepared substrates (with the exception of the acid-etched SiC, where there was essentially no difference in wear lives). This generally shorter wear life of HT films could relate to the removal of active sites at the deposition temperature of 220 °C.Finally, a comparison of the HT films on the methanolrinsed and acid-etched SiC reveals a much lower initial COF for the acidetched surface. Only HT films were observed to have parallel orientation and this orientation has been shown to depend on surface pretreatment (Fig. 1). Therefore the low initial COF for the HT film on etched SiC supports the notion that films with greater parallel orientation of crystallites have lower initial friction. A somewhat surprising result of this work was the relatively poor performance of the MoS2 films deposited on the molybdenum interlayer surface. In comparison with the steel substrates, the molybdenum interlayer films had much shorter wear lives (Table 1), indicating that simply having metal atoms at the interface is not sufficient to improve wear life. The chemical composition of the Mo/SiC interface was examined with an XPS sputter profile. The atomic concentrations of all species observed are given as a function of sputtering time (Fig. 3). The profile indicates that a significant amount of hydrocarbon and oxygen contamination exists on the substrate surface and that a small amount of silicon has migrated to the top

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SPUTTER TIME, mm Fig. 3. XPS sputter profile through a 10 nm molybdenum film evaporated onto an HP SiC substrate. The percentage composition values were obtained using the Surface Science Instruments (SSI) sensitivity factors.

of the molybdenum layer. After brief argon ion sputtering, the surface

contaminants and silicon were removed. The “bulk” molybdenum film is observed to have approximately 10% carbon (as carbide with an XPS binding energy (BE) of 283.1 eV) and a trace of oxygen (5%—7%). The carbide signal must result from the molybdenum film and not from the SiC substrate, because the silicon substrate signal is 5—10 times weaker than the carbide signal. Near the Mo/SiC interface, several forms of silicon are detected by XPS, as shown in the Si 2p data compared with clear SiC in Fig. 4. The best fits to the interface data (right side ofFig. 4) required three spin—orbit split doublets. The most intense doublet, with a 2P3,2 BE of 100.4 eV, is ionization from SiC. The 2P3/2 feature at higher BE (102 eV) is evidence for silicon oxides, which were probably present before molybdenum deposition. The low BE 2P3/2 component at 99.1 eV indicates that an interfacial reaction has occurred, with electron density shifted to silicon atoms. The BE is slightly lower than expected for elemental silicon (99.4 eV) and probably indicates the formation of a molybdenum silicide phase at the interface. (A similar peak is not seen when clean SiC is ion bombarded under the same conditions.) This reaction should also free interface carbon atoms, permitting them to diffuse into the molybdenum film and form molybdenum carbides. In addition, a similar film deposited on a tantalum substrate had significantly less carbon (less than 4%), eliminating residual gases as the only carbon source. The sputter profile results have several implications pertinent to the study of MoS 2 films in this work. The molybdenum film surface has species expected on any metal surface prior to r.f.-sputtered MoS2 film deposition (carbon contamination, metal carbides and metal oxides), yet MoS2 film life

659 DATA

FITS

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—

107

105

103

101

99

BINDING ENERGY, eV

97

95

107

105

I

103

101

99

97

95

BINDING ENERGY, eV

Fig. 4. Comparison of the date (left) and best fits (right) of the Si 2p peaks near the Mo/SiC interface after 29 mm of sputtering (top curves) and on a clean SiC substrate (bottom curves). The top spectrum was fitted with three spin—orbit split doublets (broken lines), as indicated on 2Pa 2Pi~ the figure, each having a constrained Si 1z— 2splitting of 0.62 eV and a 2:1 intensity ratio. The sum of the doublets is given by the full line with the raw data as points on the fits.

was not improved. Insight into how molybdenum reacts with SiC could be important in the deposition of MoS2 films. It is evident that molybdenum reacts with SiC by donating electron density to the silicon atoms at the SiC surface, forming an interfacial suicide, while freeing carbon atoms to diffuse into the molybdenum film. Therefore it seems that the molybdenum film would adhere strongly to SiC and that the failure of the MoS2 film on these samples was probably at the MoS2/Mo interface. When single-crystal SiC was exposed to both oxygen and sulfur at RT both species strongly chemisorbed, as detected by XPS. The oxygen coverage saturated after exposures of approximately 1000 L (1 L = 1 x 10-6 Torr s) at RT, as shown in Table 2, which gives the normalized relative XPS peak intensities [9]. A single 0 is peak was observed after RT exposure of SiC with an oxide-like BE of 532.0 eV and a very small, high BE shoulder. A 20% increase in the 0 is intensity occurred after the 500 °Canneal of the SiC in an oxygen ambient. Very little change was observed on both the silicon and carbon core levels after any RT exposures, although high BE features did appear after heating of SiC in the i0-~Torr oxygen ambient. Figure 5 shows the S 2p spectra following two sulfur source exposure times and a 500 °Canneal of the SiC. At all exposures, two spin—orbit split doublets are needed to fit the data, as shown on the right-hand side of the figure, indicating theis presence of two forms of sulfur. A lower BE form 2p~,~) more evident at low coverage (top spectrum), while a (162.3 eV S (approximately 163.0 eV) increases with exposure. The total higher BEfor form S 2p normalized XPS peak intensities are given in Table 2. After a 500°C anneal in UHV, the total sulfur intensity decreased by about 25%, with

660 TABLE 2 Normalized relative 0 is and S 2p XPS peak intensities. The S 2p and 0 is peak intensities have been divided by the Si 2p peak intensity and by the respective photoionization cross-sections (ax ,,~)(cross-sections were obtained from ref. 9) 0 is/(Si Oxygen exposure iOOLO2 300L02 1000L02 300°C,5 x iO—~Torr 7Torr 500°C,5 x lO Sulfur exposure a 30s 90s 210 s 500 °C

2p)(ao 1~)

S 2p/(Si 2p)(a 5 2p)

3.5 4.8 5.8 5.8 7.0

0 0 0 0

1.0 1.8 1.5 1.0

5.5 10.5 14.5 10.0

Oxygen predoee 500°C,5 x iO~02 150 s S exposure 500°C

5.0 5.3 5.5

0 9.1 6.4

Sulfur predose 210s S exposure 500 °C 300°C,5 x i0~ 02

1.5 1.0 1.3

14.5 10.0 7.3

~A small amount of oxygen was always detected after the sulfur exposures and probably arose from impurities in the sulfur source. The BE of the 0 is peak for this oxygen species was very similar to that observed for oxygen exposures. There was no high BE form (greater than 165 eV) of sulfur, which would indicate an interaction of the sulfur with oxygen.

virtually all signal loss occurring in the higher BE feature. These results indicate that the lower BE form is strongly bonded to SiC and that some charge transfer has occurred during formation of this bond, making the sulfur more sulfide-like and causing the XPS chemical shift to lower BE [10]. The higher binding energy form (163.0 eV) is closer to the binding energy of elemental sulfur, indicating less charge transfer from SiC, resulting in a more weakly bound sulfur phase. A final important result was observed in the competition of oxygen and sulfur for surface bonding sites. Experiments were conducted by adsorbing oxygen or sulfur on the surface, followed by exposure of the surface to the other chalcogen. The results of the experiments are summarized in Table 2. Preadsorption of oxygen inhibits the amount of strong sulfur adsorption by approximately 35%, and preadsorption of sulfur essentially prohibits subsequent adsorption of oxygen under the conditions studied. These results suggest that oxygen and sulfur compete for strong adsorption sites.

661

DATA

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(b)

~

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171

.

I I I

Z~c~

(C)

I I I Ii-’-~L I I Ir-m’~-~L 159 171 159 BINDING ENERGY, eV

Fig. 5. Comparison of the data and best fits of the S 2p peaks for the following conditions on a single-crystal SiC surface: (a) after a 30 s sulfur exposure at RT; (b) after a 210 s sulfur exposure at RT; (c) after a 500 °C UHV anneal of the surface in (b). The spectra were fitted with two spin-orbit split doublets with a constrained 2p 3~2—2p112splitting of 1.2 eV and a 2:1 intensity ratio. The doublets are indicated by the broken lines with their sum given by the full line.

4. Conclusions This work has shown that the pretreatment of SiC surfaces prior to r.f. sputter deposition of MoS2 films influences the tribological properties of the lubricating films. Films deposited on acid-etched SiC had significantly shorter wear lives than films deposited on methanol-rinsed substrates. In previous studies, the methanol-rinsed substrates were shown to have much more carbon-based contamination than the acid-etched SiC, and it appears that these carbon (or C—0) sites act as active sites for the growth of edge-oriented films with greater adhesion. In addition, HT films deposited on acid-etched SiC had measurably lower initial COFs, consistent with the enhanced amount of basal-oriented MoS2 crystallites in these films. Even though active sites were removed from the acid-etched substrates, the MoS2 films were still being deposited on contaminated SiC surfaces, and in all cases, the performance of the films depended on how strongly the MoS2 films bonded to the contaminant layers and how strongly the contaminant layers bonded with the SiC. The chemical reactivity of SiC surfaces with important species in MoS2 film growth (molybdenum, sulfur and oxygen) was studied using XPS to determine how strongly SiC bonded with these species. In these experiments, molybdenum reacted with SiC, forming an interfacial silicide and freeing carbon to migrate into the molybdenum film. During exposure of atomically clean SiC to both sulfur and oxygen, monolayer amounts of these species

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bonded very strongly to SiC (remaining on the surface following 500 °C anneals) and competed with each other for strong surface bonding sites. In summary, clean SiC should bond strongly with the individual components normally found in r.f. sputter deposited MoS2 films. This result indicates that highly adhesive MoS2 films could probably be produced in a clean deposition environment on an SiC substrate free of carbon and oxygen contamination.

Acknowledgments We thank R. Bauer of The Aerospace Corporation for preparing the MoS2 films and M. Tueling of Aerospace for making the molybdenum films. P. Davis, G. Meldrum and B. Buller at Hughes provided invaluable assistance in the tribotester wear tests. The work was funded both at Aerospace and Hughes by the Defense Advanced Research Projects Agency (DARPA).

References 1 2 3 4 5 6 7 8 9 10

P. D. Fleischauer and R. Bauer, Tribol. Trans., 31 (1988) 239. P. D. Fleischauer, Thin Solid Films, 154 (1987) 309. M. R. Hilton and P. D. Fleischauer, Mater. Res. Soc. Symp. Proc., 140(1989) 227. S. V. Didziulis and P. D. Fleischauer, Langmuir, 6(1990) 621. P. A. Bertrand, J. Mater. Res., 4 (1989) 180. J. R. Lince and P. D. Fleischauer, J. Mater. Res., 2 (1987) 827. W. Heegeman, K. H. Meister and H. Kayek, Surf. Sri., 49(1975) 161. M. N. Gardos, H.-S. Hong and W. 0. Winer, Tribol. Trans., 33 (1990) 209. J. J. Yeh and I. Lindau, At. Data Nuci. Data Tables, 32(1985) 1. C. D. Wagner, W. M. Ross, L. E. Davis, J. F. Moulder and G. E. Muilenberg (ads.), Handbook of X-ray Photoelectron Spectroscopy, Perkin—Elmer, Eden Prairie, MN, 1979.