754
Anal. Chem. 1981, 53, 754-758
Surface Analysis of Raney Nickel Alloys Joseph C. Kleln and David M. Hercules* Department of Chemistry, University of Pittsburgh, Pittsburgh, Pennsylvania
The surfaces of freshly polished Ni/AI (Raney nickel) alloys: N13AI, NIAI, Ni,A13, and NiAIS4- eutectic were studled by use of electron spectroscopy for chemlcal analysis (ESCA), Auger electron spectroscopy (AES), and secondary ion mass spectrometry (SIMS). I n all cases alumlnum was observed to segregate to the surface, the degree of segregation increasing with the bulk concentration of AI. The segregation can be predicted from the solid-liquid phase diagram of the Ni/AI system. Surface free energy calculations for substitutional pure blnary alloys gave qualitatlve agreement with empirical results for the Ni/AI system. The lack of quantitative agreement between theory and experiment is attributed to formatlon of Intermetallic compounds and oxidation effects. NaOH preferentially leached the NlA13 t eutectic and Ni2AI,.
Raney nickel catalysts display high catalytic activity for hydrogenation reactions. Variations in the preparation and composition of the alloys are known to cause variation in the activity of these catalysts. Even though numerous studies on preparation and catalytic activity exist, an omitted area concerns the surface characterization of Raney alloys. A Raney nickel catalyst begins as an Ni/Al alloy that is leached with concentrated sodium hydroxide to remove the aluminum to obtain a highly dispersed nickel catalyst. Unpublished results of Okamoto (1) indicate that hydrogenation activity changes with surface composition for Raney nickel catalysts. It has also been shown that the active catalysts could contain three of the four possible intermetallic compounds (NiA13, Ni,A13, NiAl, Ni&) (2)of the nickel-aluminum alloy system (3) (Figure 1). Furthermore, the nickel-aluminum intermetallics exhibit different leaching rates with concentrated sodium hydroxide (4). Of the four stoichiometric compounds, active Raney catalysts are produced only by leaching Ni2A13 or NiA13 (5). To investigate the nature of the differences in catalytic activity, we examined the surface properties of the Ni/Al alloys as a function of bulk composition and chemical leaching. Electron spectroscopy for chemical analysis (ESCA),Auger electron spectroscopy (AES), and secondary ion mass spectrometry (SIMS) were used in this study to monitor the surface composition and chemical states of the alloys. ESCA is a surface-sensitive (five atomic layers) technique that is quantitative and has the ability to distinguish between different chemical states (6). AES also is surface sensitive and quantitative with the ability to obtain depth profiles when coupled with ion sputtering (7). SIMS involves sputtering of a surface with kilovolt energy ions, affording mass analysis of both atomic and molecular fragments. When peak intensities are corrected for relative ion yields and transmission characteristics of the mass filter, semiquantitative SIMS results can be obtained (8).
EXPERIMENTAL SECTION ESCA spectra were obtained by the use of an AEI ES200 electron spectrometer with a DSlOO data system. An aluminum anode (A1 Ka 1486.6 eV) was operated at 12 kV and 22 mA. The base pressure was below 2.0 X lo4 torr. The digital data obtained from the DSlOO data system were processed with an HP 2114A 0003-2700/81/0353-0754$01.25/0
15260
computer. Overlapping peaks were deconvoluted with a modified nonlinear least-squares fitting routine along with a DuPont 310 curve resolver. Binding energies were measured with a precision of f0.15 eV referenced to the C Is line at 284.6 eV. Auger spectra were obtained with a Physical Electronics Model 545 Auger electron spectrometer using a 5-keV primary electron torr. For beam. The operating pressure was below 5.0 x depth profiles, 5.0 X torr of argon was used with an ion beam voltage of 5 keV and an emission current of 30 mA. Peak-to-peak heights of the Ni LMM and the A1 KLL lines were measured to calculate relative intensities. Relative surface concentrationswere calculated by using published elemental sensitivity factors (9). A 3M Model 610 secondary ion mass spectrometer having a base pressure below 5.0 X lo4 torr was used to obtain SIMS spectra of the alloys. 20Newas used as the sputtering gas. The instrument was backfilled to 5.0 X lo4 torr and operated at a beam voltage of 2 keV. The peak intensities were measured for the Alt (27 m / e ) and the Nit (58,60 m / e ) peaks and corrected for quadrupole transmission (10) and secondary ion yields (11) to obtain a semiquantitative analysis of the alloy surfaces; these results were compared with the ESCA and AES results. Alloys were prepared from nickel shot (99.99% pure) obtained from Alfa Ventron and aluminum blocks (99.99% pure) obtained from the University of Pittsburgh School of Engineering. Each component was stoichiometrically measured and the mixture melted in an argon atmosphere arc furnace. The alloys were annealed under argon. The samples were weighed before and after melting to ensure no mass loss occurred during the melting process. The following intermetallic compounds were confirmed by X-ray diffraction: Ni3Al,NiAl, Ni2A13,and NiA13 + eutectic. The alloys were cut with a diamond saw into appropriate flat samples which were polished with a silicon-base sand paper using deionized water as a lubricant. Samples were immediately placed into the spectrometers for initial analysis and were then leached for 1 h with 20% NaOH solution at ambient temperature (-20 OC). After leaching, the samples were washed with deionized water and dried under vacuum. For AES analysis the leached samples were transferred in air immediately into the instrument. For ESCA analysis, the alloys were leached under an inert (N2) atmosphere and, after evacuation,transferred to the spectrometer in a sealed probe (12) to prevent oxidation.
RESULTS AND DISCUSSION The bulk properties of the Ni/A1 (Raney nickel) alloy system change with the concentrations of the elemental components. For example, this system forms four different intermetallic compounds, each with a unique crystal structure. Since the bulk features (i.e., crystal structure, coordination number, and nearest neighbors) change with composition, the surface properties of these intermetallics are also expected to change. The surface analysis results suggest this trend. The Ni 2p3,, photoelectron peak was examined to determine the chemical state of Ni on the surface of each polished alloy sample (Figure 2). For all samples a signal was observed at a binding energy of 852.2 f 0.15 eV which can be assigned to metallic nickel (13). In the case of NiA13, a signal due to Ni2+was also observed at 856.3 f 0.15 eV (13). Its presence can be attributed to oxidation of the free nickel in the eutectic which resulted from the inability to obtain a pure NiAl3 phase (14). A weak peak detected adjacent to the metallic nickel peak (-855.9 eV) for the NiAl sample was assigned as a satellite peak, probably due to a shake-up process (15). The chemical state of the A1 was determined by examining the A1 2p photoelectron line. Two A1 peaks were observed 0 1981 American Chemical Society
ANALYTICAL CHEMISTRY, VOL. 53, NO. 6, MAY 1981 WE:GHT
755
PERCENT N:
Binding Energy ( e V )
Figure 3. AI 2p and Ni 3p ESCA spectra for freshly polished intermetallic compounds of the NVAI alloy system.
xN I Figure 1. Phase diagram of the Ni/AI system. Reprinted with permission from Hansert et al. (1958). Copyright 1958 McGraw-Hill.
Atomic %
NI
in Bulk
Figure 4. Relative intensitiesof the Ni and AI ESCA signals for pollshed NI/AI alloys.
- 868
858
848
Bindinq Energy ( e v )
Figure 2. Ni 2p3,, EISCA spectra for freshly pollished intermetallic compounds of the Ni/AI alloy system.
a t binding energies of 74.3 f 0.15 eV and 71.9 A 0.15 eV in each of the polished alloys (Figure 3). They can be assigned to A1(3+) and A1(0:),respectively. Examination of the Ni 2pljZand A1 2p ESCA spectra for the polished alloys (Figures 2 and 3) indicates the aluminum is preferentially oxidized over nickel on the surface. The preferential oxidation of aluminum correlates with a higher heat of formation for aluminum oxide than for nickel oxide (16).This result agrees with previous Ni/Al alloy surface work by Storp et al. (27). They analyzed a 50% by weight Ni alloy which contained mixed phases of NiA13 and Ni2A13. Storp et al. also concluded that aluminum was preferentially oxidized on the surface. There is some overlap between the A1 2p and Ni 3p signals as can be seen in Figure 3, so it was necessary to deconvolute
the spectra to obtain accurate relative intensities. The A1 2p and the Ni 3p (66.3 f 0.15 eV) signals were used for quantitative measurements because the kinetic energies of the photoelectrons are sufficiently similar to ensure the same sampling depth. The relative peak areas of the A1 and Ni peaks are plotted as a function of bulk metal concentration in Figure 4. Figure 4 shows the deconvoluted results of the Al2p and Ni 2p ESCA spectra plotted as the relative intensity of each signal to the total signal. The surface concentration of metallic aluminum does not change significantly among the four alloys as a function of composition. However, the aluminum oxide signal decreases linearly with increasing bulk Ni concentration for the Ni-poor alloys, with a discontinuity occurring a t 50% Ni. The Ni signal increases linearly with increasing bulk Ni content and also shows a change in slope at 50% Ni. The observed discontinuityindicates a surface change as the Ni/Al alloy system goes from aluminum-rich to nickel-rich bulk composition. The degree of aluminum oxidation, however, varies as a function of bulk composition. A linear relationship exists between the degree of oxidation of the aluminum and the bulk composition. As the aluminum increases in the bulk, so does the extent of oxidation of the aluminum. The degree of oxidation will vary with the amount of free aluminum available to be oxidized. This free aluminum appears through migration from the bulk. Therefore, the extent of oxidation is an indication of the extent of segregation of A1 to the surface. Surface segregation in this case is the enrichment of one component of a binary alloy (Al)in the upper few atomic layers (ca. 5 atomic layers) relative to the bulk of the alloy. This
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ANALYTICAL CHEMISTRY, VOL. 53, NO. 6, MAY 1981
Table I. Determination of Ni Surface Concentration to the Ni Bulk Concentration Ratios by AES alloy NiAl, Ni,Al, NiAl Ni.Al
20
I
I
I
I
40
60
80
IC0
Ni surface concn/Ni bulk concn 0.54 0.58
0.71 0.19
Atomic 'lo N i in Bulk
Flgure 5. Relative Ni intensities of static SIMS and AES for polished Ni/AI alloys: AES, ---; static SIMS, -.
0 E e o m On
IO
20
30 4 0 Time ( m i n )
50
Flgure 7. Auger depth profile of Ni/Ai. Atomic %
Ni in B u l k
Figure 6. Relative concentrations of Ni on the surfaces of polished NVAI alloys.
criterion is employed because the instrumental techniques employed have this approximate sampling depth. In order to substantiate the ESCA results, other surface techniques were employed. Because static SIMS has a Sampling depth of one or two atomic layers and AES has a Sampling depth of five to ten atomic layers which is essentially the same as that of ESCA, the information obtained should be complementary with the analytical information derived from ESCA. By measurement of the intensities of Ni and A1 signals in static SIMS and AES, the relative Ni signal vs. bulk concentration could be examined in the same manner as in Figure 4. The static SIMS A1 (27 m / e ) and Ni (58,60 m / e ) intensities of the polished alloys were corrected for quadrupole transmission (10) and relative ion yields (11)by using published values. The fractional corrected Ni intensity of the combined Ni and A1 signals indicates the same surface change (Figure 5) as was observed by ESCA. In the same manner AES was used to examine the polished alloys. The relative intensity of Ni, determined by measuring the peak-to-peak height of the Ni LMM and the A1 KLL Auger lines, indicated the same discontinuity of the relative Ni intensity (Figure 5 ) . By application of published elemental sensitivity factors (9) to the Auger results, the relative surface concentrations of Ni for the alloys were determined. These are plotted as a function of bulk Ni concentration in Figure 6. Because the relative Ni concentrations lie below the 45O diagonal line, it can be concluded that A1 predominates on the surface, i.e., A1 surface segregates. By calculation of the Ni surface concentration/Ni bulk concentration ratios from the Auger data (Table I), the degree of surface segregation can be estimated as a function of bulk concentration. The nickel surface concentration to nickel bulk concentration ratio (Table I) is evidence that the degree of segregation is changing as a function of bulk composition. For instance, the NiA13surface contains about 50% less nickel than the bulk, while the Ni3A1surface
is only 20% deficient in nickel. These results are in agreement with the lower surface free energy of aluminum affording its segregation to the surface (18). In view of the ESCA, AES,and SIMS relative intensity data (Figures 4 and 5), the surface is changing in character as the bulk changes from predominantlyaluminum to predominantly nickel. This is concluded because the slope of the relative nickel intensity changes significantly at about 50 atomic percent nickel in the bulk. Figure 4 shows that the amount of metallic aluminum on the surface of an alloy does not change significantly with bulk composition. However, the oxidized aluminum displays the same abrupt change as the nickel. In light of this, one can conclude that the oxidation and segregationof the aluminum change as the bulk changes from predominantly nickel to aluminum. Therefore, the increase in surface segregation and reactivity of NiAl, and Ni2A13 infers that the chemical stability of these phases is less than that of the nickel-rich alloys. Since differing degrees of A1 segregation are observed on the surfaces of the polished alloys, it is expected that the oxide thickness should vary with bulk Ni concentration. Relative oxide thicknesses were measured by Auger depth profiling. A typical result is shown in Figure 7 . For depth profiles, the peak-to-peak heights of the Ni LMM, A1 KLL, 0 KLL, and C KLL Auger peaks are plotted as a function of sputtering time. In Figure 7, the small increase in oxygen signal just after the ion beam was turned on was caused by initial sputtering of carbon contamination on the surface. In depth profiles, the sputtering time is assumed to be linearly related to the thickness of the sputtered layer. By measurement of the time needed to sputter to the interface between the oxide and the underlying element (i.e,, Ni), an estimate of the oxide thickness for different alloys can be obtained (19). A linear relationship exists between the oxide thickness and the bulk concentration for the Ni/A1 system based on AES depth profiles. It is observed that the oxide thickness decreases with increasing bulk Ni concentration. In order to obtain an estimate of the actual oxide thickness, aluminUrri samples having known oxide thickness were sputtered under the same conditions as the Ni/Al alloys, to correlate sputtering times to oxide thickness. Aluminum oxide films were employed because the oxide films
ANALYTICAL CHEMISTRY, VOL. 53, NO. 6, MAY 1981
: I .-
757
\
Alornic %
NI
in BiJlh
Flgure 8. NaOH leaching results of Ni/AI alloys.
of the alloys are mainly aluminum oxide. The estimation of the oxide thicknesses is as follows: NiA13 + eutectic, 66 A; Ni2A13,50 A; NiAl, 40 A; Ni3Al, 18 A (ij min of sputtering corresponds to approximately 35 A). This result strongly suggests that differing amounts of AI are migrating to the surface and are being oxidized as the Ni content of the alloy is changed. The Auger depth profiles (Figure 7) illustrated that the thickness of the oxide layer varied as a function of bulk composition. The Ni and 0 Auger signals were used to determine changes in oxide thickness because each displays a significant concentration gradient between the oxide layer and the bulk. The oxygen changes from a high surface concentration to nearly zero in the bulk while the nickel increases from a low surface amount to its bulk concentration. Casual inspection of the depth profile (Figure 7) could infer a misleading conclusion that aluminum does not display a significant change in relative concentration from the surface to the bulk. However, when the aluminum migrates to the surface and is oxidized, oxygen dilutes the relative aluminum concentration enough so that only a small change in Auger signal intensity is observed. Also, a noticeable depletion of aluminum a t the beginning of the bulk interface could be expected because of possible %voids left behind by the migrating Al atoms. However, this is not observed in the depth profiles because migration occurs as a continuum from the bulk to the surface where vacancies made by the migrating atoms are filled by the atoms adjacent to them. The variation of the oxide thickness agrees with the ESCA results for the polished samples. Again the increase in oxide thickness as the bulk aluminum increases displays the increasie in reactivity and segregation of the aluminum-rich alloys. T o investigate a more drastic treatment of the alloys, we employed chemical leaching with NaOH to determine the degree of reactivity of the aluminum. After the sample was leached for 1h in 20% NaOH at ambient temperature, the alloy surfaces were analyzed by using both AES and ESCA. The difference between the relative Ni Auger intensities of the leached and freshly polished alloys (INi/(INi Id)leached - I N ~ / ( I N+~ I N )polished) plotted as a function of bulk composition is shown in Figure 8. The NiA13 and the Ni2A13 samples showed significant Ni enrichment on the surface after leaching. However, little if any Ni surface enrichment was observed for NiAll and Ni3Al. By comparison of the ESCA A1 2p and Ni 3p spectra of the leached alloys (Figure 9) with the same spectra of the polished alloys (Figure 3), enhancement of the Ni signals relative to the A1 signal after leaching for NiA13 and Ni2A13 is clearly evident. The Ni 3p binding energy (66.4 f .15 eV) for the leached alloys indicates that the Ni remains in the metallic state during the base leaching
+
Figure 0. AI 2p and Ni 3p ESCA spectra of leached alloys.
process, This is also supported by the Ni 2p1p ESCA line of the leached alloy which lies at 852.3 f 0.15 eV. The A1 2p ESCA spectrum after leaching indicates that additional oxidation of A1 occurred for the Ni3A1. However, for the other alloys the extent of oxidation of the aluminum remained about the same, The NiAl and Ni3Al alloys indicate very little if any reaction with base; i.e., they are inert. However, the Ni2A13 and NiAls react significantly with base. These results are consistent with the air oxidation results (Figure 4). Also, the differing degrees of leaching agree with previous Ni/Al alloy leaching work (4). The surface analysis of the polished and leached Ni/A1 alloys shows NiAl and Ni3A1intermetallics are more resistant to surface segregation and chemical attack. On the basis of the surface analysis, it appears that the aluminum-rich intermetallic compounds are less chemically stable than the nickel-rich compounds in view of their greater surface segregation and oxidation. This is clearly shown in the greater depletion of nickel on the surface in the aluminum-rich alloys and the larger extent of oxidation of A1 determined by ESCA and AES. Also, under drastic chemical attack the aluminum of NiA1, and Ni2Al8is oxidized and removed. This was observed by a greater amount of nickel on the surface after leaching. Also, the NiAl Al2p and Ni 3p ESCA spectra before and after leaching indicate very little change because the relative intensities of the signals are essentially the same. However, Ni3A1did display some additional oxidation of the aluminum after leaching. The aluminum metal and oxide that remained on the surface of Ni2AI3and NiA13 are in about the same proportion as they were on the surface of the polished alloys. Since surface segregation of the Ni/A1 alloy system is observed, it would seem appropriate to compare the surface data of this system with surface segregation models for binary alloys. Seah has proposed a model to calculate the surface energy of a substitutionally pure binary alloy system, using the equation (20) EA
= -In BASRT
(1)
where EA = surface free energy of A , R = gas constant, T =
758
ANALYTICAL CHEMISTRY, VOL. 53, NO. 6, MAY 1981
Table 11. Theoretical and Experimental Surface Free Energies substitutional pure empirically binary alloy surface determined surface free energy free energies calcd calculation from from eq 3 using eq 1 and 2 AES data alloy Ni,AI
NiAl Ni,Al,
NiA1,
ENir
EA19
h i ,
EAI,
kJ/mol
kJ/mol
kJ/mol
kJ/mol
19.0 19.0 19.0 19.0
-30.2 -30.2 -30.2 -30.2
1.8 1.5 2.0 1.8
-1.8 -1.5 -2.0 -1.9
temperature, Ti” = melting temperature of i, H” = heat of mixing, 2 = coordination number, XF = mole fraction of i in bulk, ai = atomic size of i, and M = 1if UA > LIB and M = 0 if UB > CIA. The results of this calculation for the Ni/Al alloys studied in this work are given in Table 11. The surface free energies for the Ni/Al system also can be estimated by using an empirical approach determined by AES, using the Langmuir equation (20) -=Xi“
Xi“ e x p ( - g ) 1-xis 1-xp
(3)
where Xis = mole fraction on the surface of i, X p = mole fraction of i, T = temperature, R = gas constant, and E: = surface free energy of i. These empirically determined surface free energies are also given in Table I1 (right-hand column). Comparison of empirical and theoretical surface free energies shows a discrepancy greater than 1order of magnitude. There are two reasons for the differences. First, nickel and aluminum have a high degree of elemental interaction and form intermetallic compounds (21). Substitutionally pure binary alloys are assumed to have little metal-metal interaction and the value of the coefficient on the elemental interaction term in eq 2, 1.86, is small. Therefore, the Ni/A1 system would not be expected to follow the substitutionally pure binary alloy criteria. Second, segregation can be affected by exposure to oxygen, a phenomenon observed with many systems. For example, iron surface segregates to form an oxide layer when platinum-iron alloys are exposed to oxygen but not under an inert atmosphere (22). Hence, segregation induced by oxygen exposure may also contribute to the discrepancy between theoretically and experimentally determined surface free energies. Burton et al. (23)proposed a rule that qualitatively predicts surface segregation from solid-liquid phase diagrams. The predictions are based on the temperature range over which the solid and liquid phases coexist for a particular composition. This region lies between the pure solid and pure liquid phases. If the solid-liquid coexistence for a specific composition has a large temperature range, significant surface segregation will be expected. When this temperature range is small, little or no segregation should occur. The phase diagram for the Ni/Al alloy system (Figure 11of ref 24) reveals a large solid-liquid
coexistence region above NiA13and Ni2A13and smaller regions above NiAl and Ni3Al predicting that enhanced segregation should occur for the aluminum-richalloys. This coincides with the results from surface analysis. For all intermetallic compounds studied, aluminum segregated to the surface. The degree of segregation and surface oxidation increased with increasing bulk concentration of Al. Also, the reactivity toward base increased with increasing bulk Al. These results imply a lack of chemical stability of the Al-rich intermetallics. The A1 segregation can be predicted from the solid-liquid phase diagram. However, surface free energy calculations for substitutional pure binary alloys gave only qualitative agreement with experimental results for the Ni/Al system. The lack of quantitative agreement between theory and experiment is attributed to intermetallic compound formation and the effects of oxidation.
ACKNOWLEDGMENT We are grateful to John Gasper and the University of Pittsburgh School of Engineering for providing us with the intermetallic compounds. LITERATURE CITED (1) Okamoto, Y., Osaka University, Osaka, Japan, unpublished studies, 1974. (2) Bradley, A. J. Philos. Mag. Suppl. 7 1937. 23, 1049-1067. Vaughan, N. B.. J . Inst. Met. 1937, 61, 247-280. (3) Alexander, W. 0.; (4) Freel, J.; Pieters, W. J. M.; Anderson, R. B. J. Catal. 1970, 16, 281-291. (5) Sassoulas, R.; Trambouze, Y. Bull. SOC. Chim. f r . 1964, 5, 985-988. (6) Kim, K. S.; Winograd, N. Surf. Sci. 1974, 43, 825-643. (7) Joshl, A.; Davis, L. E.; Palmberg, P. W. “Methods of Surface Analysis”: Czanderna, A. W., Ed.; Elsevier: New York, 1975;Chapter 5. (8) McHugh, J. A. “Methods of Surface Analysis”; Czanderna, A. W., Ed.; Elsevier: New York, 1975;Chapter 6. (9) Davis, L. E.; MacDonald, N. C.; Palmberg, P. W.; Rlach, G. E.; Weber, R. E. “Handbook of Auger Electron Spectroscopy”, 2nd ed.; Physical Electronics Industries, Inc.: Eden Prairie, MN, 1976;Chapter 1. (IO) UTI, Model lOOC, Operating and Service Manual, Sunnyvale, CA,
1975. (11)Cherlpln, V. T. Adv. Mass Spectrom. 1975, 7A,776-783. (12) Patterson, T. A.; Carver, J. C.; Leyden, D. E.; Hercules, D. M. J. Phys. Chem. 1976, 80, 1700-1708. (13) Kim, K. S.; Davis, R. E. J. Electron Spectrosc. Relat. Phenom. 1972, 1, 251-258. (14) Holm, R.; Storp, S. J. Electron Spectrosc. Relat. Phenom. 1976, 8 , 139-147. (15) Andrews, P. T.; Collins, T.; Johnson, C. E.; Weightman, P. J. Nectron Spectrosc. Related Phenom. 1879, 15, 39-42. (16) Mledema, A. R. 2. Metallkd. 1976, 69, 455. (17) Storp, S.;Bureshelm, K.; Wllmers, M. SIA Surf. Interface Anal. 1979, 1, 96-99. (18) Bouwman, P.; Joneman, L. H.; Holscher, A. A. Surf. Sci. 1973, 35, 8-33. (19) Castle, J. E..; Hazell, L. B., J. Electron Spectrosc. Relat. Phenom. 1977, 12, 195-202. (20) Seah, M. P. J. Catal. 1979, 57,450-457. (21) Hultgren, R.; Desai, P. A,; Hawklns, D. T.; Glelser, M.; Kelly, K. K. “Selected Values of the Thermodynamic Properties of Binary Alloys”; American Society for Metals: Metals Park, OH, 1973; pp 191-195. (22) Bartholomew, C. H.; Boudart, M. J. Catal. 1973, 29, 278-291. (23) Burton, J. J.; Machlln, E. S. Phys. Rev. Lett. 1976, 37, 1433-1438. (24) Hansen, M.; Anderko, K. “Constitution of Binary Alloys”, 2nd ed.; McGraw-HiII: New York, 1958;p 119.
RECEIVED for review October 2,1980.Accepted January 27, 1981. This research was supported by the National Science Foundation under Grant CHE78-00876.
