- Email: [email protected]
Surface compositions of copper-silicon alloys
Applications of Surface Science 14 (1982—83) 359—374 North-Holland Publishing Company
SURFACE COMPOSITIONS OF COPPER-SILICON ALLOYS Timothy C. FRANK and John L. FALCONER Department of Chemical Engineering Campus Box 424, University of Colorado, Boulder, Colorado 80309, USA Received 3 September 1982; accepted for publication 13 December 1982
Auger spectroscopy of copper—silicon alloys showed that silicon is enriched at the surface of both an ordered compound, Cu 3Si, and a solid-solution alloy (5 at% Si). The surface compositions are estimated using elemental standards. Silicon segregation is predicted by thermodynamic theories based on surface tensions of pure components and by the empirical Burton and MachIm rule. The Si(92 eV) Auger peak splits into two peaks upon alloying, making identification of alloyed silicon possible. Surface silicon is in an alloyed form for both Cu3Si and Cu0 95Si005. Above 670 K, additional silicon segregates to the surface of Cu3Si and both alloyed and elemental silicon are present. Alloying copper with silicon results in weakened silicon bonds and clean Cu3Si oxidizes much more rapidly in air than does elemental silicon. This weakened bonding may be one reason that copper catalyzes the reaction between methyl chloride and silicon to form methylchlorosilanes.
I. Introduction In the commercial production of silicone polymers, organohalosilanes are hydrolized to form silanols, which condense to form silicone polymers and water. The principal industrial synthesis of these silanes is the direct-synthesis reaction between gaseous organic halides and elemental silicon. The most important reaction in this class of solid-catalyzed, gas—solid reactions is the formation of dimethyldichlorosilane ((CH3 ) 2SiCI2) using a copper catalyst. Cu
2 CH3C1
+
Si
—*
(CH3)2SiCI2.
(1)
Several studies have demonstrated that catalytic activity for this reaction is correlated with the amount of ordered-compound Cu3Si (~phase) present [1,21,but the reaction is poorly understood. As a first step toward understanding the reaction mechanism, we measured the surface compositions of clean ~ alloys as a function of temperature using Auger spectroscopy. Two Cu3Si samples, prepared separately, were used to determine the accuracy of the Auger measurements. A dilute, solid-solution 0378-5963/83/0000—0000/$03.OO © 1983 North-Holland
360
T C. Frank, J. L. Falconer
/
Surface compositions of Cu
—
Si alloys
alloy (5 at% Si in Cu) was also compared to Cu3Si to determine any surface differences that might influence the direct-synthesis reaction. Auger spectroscopy should yield a good estimate of surface composition for copper—silicon alloys since both copper and silicon have Auger peaks in the energy region which corresponds to minimum escape depth [3]. The silicon Auger peak at 92 eV split into two peaks upon alloying. This split peak serves as a good indication of the presence of alloyed silicon. However, it also complicates surface analysis, and surface compositions were estimated using the high-energy silicon peak and the copper peaks. The enrichment of silicon on the surfaces of these alloys was compared to theory predictions.
2. Experimental Surface analysis was carried out in an ion-pumped UHV chamber equipped with a PHI CMA Auger spectrometer, a UTI quadrupole mass spectrometer, and a Varian ion bombardment gun. After bakeout, pressures of 7 x i0~Pa (5 x l0~b0Torr) were easily obtained. Samples were either mounted on a PHI rotatable sample manipulator or in a transfer mechanism which allowed samples to be changed without exposing the UHV chamber to atmospheric pressure [4]. Samples were heated from behind by radiation from a 0.25 mm diameter tungsten filament. The temperature was measured with a chromel—alumel thermocouple (0.13 mm diameter wire) spot-welded to the side of the sample. X-ray analysis of the Cu3Si alloy was done using a Buerger precession camera. A Zeiss polarized-light microscope was used to observe the optical anisotropy of the samples. Two Cu3Si samples were studied. Sample No. 1 was prepared from a melt of 23 at% Si and 77 at% Cu since silicon precipitates formed at higher silicon concentrations [5]. The alloy was made from large chunks of silicon (6 mm diameter) in order to reduce the amount of Si02 present in the melt. The copper and silicon were sealed in an evacuated quartz ampoule and heated to 1300 K in a horizontal tube furnace. Vigorous agitation of the melt/solid mixture for 20 mm was required to incorporate unreacted silicon into the melt [5]. The resulting ingot was annealed at 670 K. Sample No. 2 was obtained from Union Carbide. The Cu0 95Si005 sample was also prepared by heating copper and silicon in an evacuated quartz ampoule at 1300 K in a horizontal tube furnace. The melt was agitated for 20 mm and then cooled to room temperature in 30 mm. Samples were sliced with a diamond saw and polished with Linde B abrasive (0.05 ~m Al 203). The Cu3 Si alloys were extremely brittle, so fairly thick samples (1.5 mm) were used. The Cu095Si005 sample was not brittle2and on was eacheasily face. sliced. The surface area of each sample was approximately 1 cm
T. C. Frank, J.L. Falconer
/ Surface
compositions of Cu — Si alloys
361
The silicon sample was an n-type single crystal, Si( 100), manufactured by Monsanto and containing iO-~%phosphorus. The copper was a polycrystalline foil (99.995% purity) supplied by Alfa Products. Argon used for ion bombardment was specified as 99.9995% pure (Scientific Gas Products, research grade) with less than 1 ppm H20 and 02. Carbon monoxide (Scientific Gas Products, technical grade) used for adsorption was 99% pure. All Auger spectra were recorded with a 3 keV beam voltage and a 30 .tA beam current. The modulation voltage was 6 V peak-to-peak for Si( 1619 eV) and I V peak-to-peak for all other peaks. The analyzer was calibrated using the negative peak of the 2000 eV, differentiated elastic peak.
3. Results Both the bulk and the surface of the samples were analyzed. X-ray diffraction and polarized-light microscopy were used for the bulk analysis. Auger spectroscopy, CO adsorption and scanning electron microscopy were used for surface analysis. The Auger spectra were recorded as a function of annealing temperature. Auger spectra of elemental silicon and copper were used in estimating surface compositions. 3.1. Bulk characterization
The X-ray diffraction for a small crystallite taken from the polycrystalline Cu3 Si sample No. 1 indicated an ordered structure with six-fold symmetry and hexagonal lattice constants a 0.70 nm and c 1.46 nm. The complex structure of Cu3Si is uncertain [6—9].Three temperature-dependent phases of Cu3Si are known to exist [10]; however, the active form for the direct-synthesis reaction is not known. The structure of our sample may be the mid-temperature phase (700—830 K) given by Solberg [9]; phase This is rhombohedral tR this 95.72°. corresponds R3 to with latticelattice constants aR 0.472 nm and hexagonal constants a 0.700 nm and c 0.732 nm. The six-fold symmetry for our alloy does not agree with the three-fold symmetry of Solberg’s structure but this may be due to twinning of the crystallite, since Cu 3Si has a tendency for twin formation [6,9]. The Cu3Si stoichiometry was verified on polished samples by using a polarized-light microscope. The three i~phases and the K phase (10—15 at% silicon) are the only optically-anisotropic compounds in the copper—silicon system. The K phase is not present at the conditions used to prepare our alloys [11]. The index of refraction for an optically-anisotropic material is a function of crystal orientation. When the polished Cu3Si surfaces were rotated between crossed polarizers, the intensity changed periodically, alternating between =
=
~‘
=
=
=
=
362
T. C. Frank, J.L. Falconer
/
Surface compositions of Cu — Si alloys
extinction and full illumination of the surface crystallites. Intensity does not change periodically for isotropic materials. This verifies the Cu3Si stoichiometry of our alloys. The polarized-light microscope showed that the Cu0 95Si0 ~ alloy was optically isotropic.
Si,570K
~3Si,77OK
dN(E) dE
A
Cu3SI,570K
~,5,770
40
K
60 80 100 120 Electron Energy (eV)
Fig. 1. Low-energy Auger spectra of alloys and pure elements after annealing at indicated temperatures for 30 mm.
T C. Frank, J.L. Falconer / Surface compositions of Cu — Si alloys
363
3.2. Surface characterization 3.2.1. Scanning electron micrographs
Before placement in the UHV chamber, the samples were polished to a bright mirror finish and analyzed by SEM. The only features observed in SEM were a few defects at magnifications of 17000 x After removal from the UHV chamber following repeated ion bombardment/annealing cycles, a roughened surface was observed on the Cu3Si samples. Scanning electron micrographs .
showed pitted areas 0.5 ~smwide and several hundred ~.tmlong. These pits were aligned and comprised about 10% of the total surface. 3.2.2. Surface cleaning
Cycles of argon ion bombardment and annealing were used all the 2 into6.7clean X l0~ Pa samples.Annealing Bombardment was570 done 500principal eV and 2—5 ~sA/cm were carbon and argon. was at K. atThe impurities oxygen, but nitrogen and sulfur were present on Cu 3Si and chlorine and sulfur were seen on copper. Long cleaning times reduced all contaminants below the AES detection limits on Cu3Si and copper. Small carbon and oxygen levels were still present on Cu0 95Si0 05 after 15 h of cycling. The 0(508 eV)/Cu(60 eV) and C(273 eV)/Cu(60 eV) Auger peak ratios were 0.05 and 0.02 respectively. Annealing Cu0 95Si005 to 770 K caused small concentrations of phosphorus and sulfur to segregate to the surface. Similar small carbon, oxygen and argon ratios were observed for the silicon sample. 3.2.3. Auger spectra
The pure silicon spectrum was very similar to that reported in the literature [12]; peak locations were within 1 eV. Auger peak energies for pure copper were also similar to those reported in the literature except for small differences in peak locations [12,13]. Several changes in the Auger spectra occurred on alloying, however. Fig. 1 shows the low-energy Auger spectra of copper, silicon, and the alloys. The low-energy silicon peak was significantly altered by alloying. Pure silicon has a single peak at 92 eV while both copper—silicon alloys exhibited two peaks at 90 and 94 eV. The peak at 82 eV for pure silicon shifted to 85 eV after alloying and its relative peak height increased from 0.01 to 0.15. The relative peak height is the ratio of secondary to primary peak height in a given peak ensemble (see table 1). Relative peak heights for the other silicon peaks also changed on alloying. The 1601 eV peak shifted to 1596 eV and its relative peak height decreased from 0.32 to 0.10 (see table I). The copper peaks, however, were not altered by alloying. Auger spectra of Cu3 Si were run both before and after silicon spectra, so that accurate comparisons could be made. All alloy spectra were reproducible within ±5—10% of peak height and ±1 eV in peak location. The shift in Auger peak locations are not due to inaccurate calibra-
364
T. C. Frank, J.L. Falconer
/ Surface
compositions of Cu — Si alloys
Table I Relative silicon peak heights Elemental silicon
Cu
3Si alloy
Peak location (eV)
Relative peak height ~
Peak location (eV)
92~ 82 16l9~ 1601 1582 1561 1543 1520
I 0.01 I 0.32 0.13 0.27 0.12 0.05
90,94~> 85 l620~ 1596 1579 1560 1538 1518
Relative peak height a)
0.15 0.10 0.10 0.25 0.05 0.10
~ Relative peak height; the ratio of secondary peak height to primary peak height in a given peak ensemble b) The reference peak used to calculate the relative peak heights for energies that follow in the table. ~ Si(94 CV) is used as the reference peak for the alloy.
tion; the copper peaks at 60 and 917 eV were the same for both copper and the alloys. The two silicon peaks at 90 and 94 eV changed to a single peak at 92 eV when the Cu3 Si alloy was annealed above 670 K, as shown in fig. 1. This single Si(92 eV) peak is much broader than the corresponding peak measured on pure silicon and may be a combination of spectra for alloyed and unalloyed silicon. The split silicon peak measured on Cu0 95Si005 did not change to a single peak when this alloy was heated between 300 and 830 K. Fig. 2 shows peak height as a function of annealing temperature for Cu3Si alloy No. 2. The data were obtained by annealing the ion-bombarded sample at the temperatures indicated in fig. 2, beginning at 370 K and progressing to higher temperatures. Auger spectra were recorded after sufficient annealing times that the Auger peak heights reached steady-state values: 30 mm at 370 K, 20 mm at temperatures up to 670 K, and 10 mm at higher temperatures. In separate experiments, the sample was annealed for 145 mm at 370 K, 60 mm at 470 K, 60 mm at 570 K, and 30 mm at 770 K with the same results as in fig. 2, indicating that the equilibrium surface was rapidly attained. 3.2.4. Surface composition
The ratio of an Auger peak height in the alloy to the same transition peak height in the pure component should estimate the surface concentration of that component in the alloy. Table 2 lists these ratios. With few exceptions the
T C. Frank, J. L. Falconer
I
I
I
/
Surface compositions of Cu — Si alloys
I
I
Si
90,94
5
,,—~~ci—-—a-Si92
Cu6~
~°
-
SI9~
~
~
Si~6~0~
~
Cu917
-
Si1620
-
CX 15)
0
-(X15) 2
-
~
-
~
I
Si92
-
I~
0
-
~Cu60
-
-
-
Si90,94
-
1
-
split height —i
-
0
-
300
365
I 400
500 600 Temperature (K)
I 700
I 800
-
Fig. 2. Equilibrium Auger peak heights for Cu3Si sample as a function of annealing temperature.
agreement between values obtained on samples No. 1 and No. 2 of Cu 3Si is very good. Also, for most ratios, the silicon atomic fractions estimated from the copper peaks (by subtracting the copper peak ratio from one) are in reasonable agreement with the fractions estimated from Sm( 1619 eV). Following annealing of Cu3Si at 770 K, the more surface-sensitive copper Auger peaks indicate a much higher silicon concentration than the high-energy copper or silicon peaks. However, for Cu 0.95 Si 0.05’ annealing at 770 K increased all the Auger peaks, complicating estimates of surface composition. For all samples, the Si(94 eV) peak indicates a silicon surface concentration 1.5 to 2.5 times as large as that obtained from the other peaks. Since this silicon Auger peak shape changed upon alloying, comparisons using this peak are not meaningful. To obtain better estimates of surface compositions, the method of Hall and Morabito was used to account for matrix effects [14]. This method uses a correction factor F which accounts for differences between Auger transitions in the alloy and pure silicon. This correction is for dilute alloys, evaluated in the limit of pure copper, and is independent of the copper transition used. This method is not strictly valid for concentrated alloys such as Cu3Si, but F is not
366
T. C. Frank, J.L. Falconer
/
Surface compositions of Cu — Si alloys
aV 0.
V .-
C
9
N
~
V
V
V
~o o’~~——~ ~a N — ~ ddd~~~d~ .2 C
-~
V — > V
0
N
0.
>
N-N’C~0’N’C ~0
U
~O
‘~ 0’
N 0’
‘O
~
a
a -o
‘5
0
V
‘c~ C
.9 I-
Zr.. ~
U
N s ~
N ‘.0 ‘.0
~ ‘.0
~O N en 0’
~
~ 0’
a ~~
.~
—
I ~ 0.
~0
<
0.
v
C
U0
V
VO’ — N
~
0’ 0’ ‘t~ ‘0
— ‘i”.
~n
‘0 ~O
‘
C ~
.5 —
~
t
E
‘0 ‘0
~
C N
‘0
U
CC N N r’. N
‘0
‘0 ~ ‘0
5
CC N N
0,
—
~a .2
V
~
o~oaa0~~ a
~
cc
cc
.OQCUQ.0QV
~
~
a a
— ~0 —
.0
a 0’
0
U
Z
N
0
Z
..~
N
a
0
U
U
U
.0
T. C. Frank, J.L. Falconer
/
367
Surface compositions of Cu — Si alloys
a strong function of composition [14] and the atomic densities of Cu and Cu3Si 3) [7]. The method assumes are and and 7.8 Xthat 1022 atoms/cm peaksimilar shapes(8.5 do x not1022 change surface roughness and diffraction effects are negligible. Diffraction errors should be small since Auger emission from Si(100) is nearly isotropic [15] and other samples are polycrystalline. Hall and Morabito [14] used two correlations for escape depths in correcting for matrix effects. When corrected for the 3 keV primary beam voltage used in this study, the values for F are 0.665 (Penn correlation) and 0.905 (Seah and Dench correlation) [14]. Table 3 lists the silicon atom fractions obtained with the Hall and Morabito matrix corrections for both Penn’s and Seah and Dench’s (SD) correlations for escape depth. All calculations are based on the Si(l6l9 eV) transition since the more surface-sensitive transition at 92 eV changes shape with alloying. For the ion-bombarded samples, surface composition estimates using the Penn correlation are very close to bulk values. Since previous studies have indicated that sputtering yields of copper and silicon in Cu—Si alloys are equal [16,17], the ion-bombarded sample should have a surface composition close to the bulk. Thus the Penn correlation may yield the more accurate estimates. However, it is possible that the Auger beam, which heated the samples to 350—400 K, annealed the surface and caused silicon enrichment. Table 3 shows that following annealing at 570 K, the surface compositions of Cu 3Si and Cu0 955i005 are at least 30 at% Si and 10 at% Si, respectively. Agreement between the two Cu3 Si samples is very good. Annealing Cu3 Si at
Table 3 Silicon atomic fraction a) for various copper peaks Sample
Surface treatment
Cu3Si (No. 1) Cu3Si (No. 2)
Cu0 95Si005
a)
Ion bombarded Annealed at 570 Ion bombarded Annealed at 570 Annealed at 770 Ion bombarded Annealed at 570 Annealed at 770
K K K K K
Cu(60 eV)
Cu(105 eV)
Cu(917 eV)
Penn
SD
Penn
SD
Penn
SD
0.24 0.30 0.29 0.33 0.50 0.07 0.10 0.10
0.30 0.37 0.36 0.40 0.58 0.09 0.13 0.13
0.23 0.29 0.29 0.32 0.45 0.06 0.10 0.10
0.29 0.36 0.36 0.39 0.53 0.08 0.12 0.13
0.25 0.29 0.29 0.30 0.37 0.06 0.10 0.10
0.31 0.36 0.36 0.37 0.44 0.08 0.13 0.13
(paIIoY/pSi)F S~
(paIloY/pca)+
(p;11oY/p~)F
The Si(1619 eV) peak was used for all calculations.
368
T. C. Frank, J.L. Falconer / Surface compositions of Cu
—
Si alloys
770 K produced a surface containing more than 40 at% Si, with higher values indicated for the more surface-sensitive peaks. Annealing Cu0 95Si005 at 770 K caused no additional change in the surface composition though both copper and silicon peak heights increased (table 2). The silicon surface enrichment observed is not an artifact due to preferential sputtering during ion bombardment. Following bombardment, the surface composition was approximately the same as the bulk and enrichment only occurred on annealing. Also, additional bombardment yielded the same bulk composition. Moreover, no preferential sputtering was detected in previous studies of Cu—Si alloys [16,17]. 3.3. CO adsorption
Preliminary CO desorption experiments were carried out for Si, Cu and Cu3Si in an attempt to determine the composition of the surface monolayer. 2) for On three surfaces COK.coverage was lowadsorption ((1—2) x l0~molecules/cm 50 Lalladsorption at 310 No additional was observed at higher exposures. Carbon monoxide adsorption on nickel was used for calibration [18].On silicon and copper, a small CO peak was observed at 370 K. A similar
/1
/1
Contaminated Cu 3Si
dN(E)
Coppe~0PP&
273
Clean Cu3Si
~
Copp~~~pper
Carbon
60 67
67 79 85
Silicon
85 94
Silicon
508
90
Oxygen //
Electron Energy (eV) Fig. 3. Low-energy Auger spectra of contaminated and clean Cu3Si sample. The clean sample was annealed at 570 K.
T. C. Frank, fL. Falconer
/ Surface
compositions of Cu — Si alloys
369
peak has been observed on Si(100) [19].Carbon monoxide is weakely adsorbed on copper [20]. A broad peak appeared at 330 K on Cu3Si for samples annealed at 570 or 770 K. The CO desorption data could not be used to determine the surface monolayer composition because CO coverage was too low. The surface could have changed significantly without affecting CO desorption. 3.4. Cu3 Si oxidation
The ion-bombarded Cu3Si alloy oxidized much more readily than Cu0 95Si005, the pure components, or the Cu3Si samples before ion bombardment. The cleaned Cu3Si samples formed a black oxide surface layer after less than 24 h exposure to room-temperature air, while the pure components and the Cuo 95 0.05 alloy showed no visible corrosion even after several months exposure. Fig. 3 shows the Auger spectra of a clean Cu3Si alloy which was exposed to methyl chloride and oxygen and then heated. Oxygen and carbon were the only contaminants observed. The intensity of the peak at 67, 79 and 85 eV increased for the contaminated surface. This plus the oxygen peaks at 475, 490 and 508 eV may indicate the presence of Si02. The relative amplitudes of the 90 and 94 eV peaks also changed but both peaks were present, indicating surface silicon was still alloyed with copper.
4. Discussion 4.1. Silicon bonding in Cu—Si alloys
Copper significantly alters silicon bonding upon alloying, as evidenced by the split Si(92 eV) Auger peak, the changes in other silicon peaks, and the increased oxidation rate of Cu3Si. Silicon bonding appears to be weakened in the alloy, particularly in Cu3Si, and this may be one reason why copper, in Cu3 Si, is a good catalyst for the direct-synthesis reaction. The weakened silicon bond in the alloys is more easily broken to form gas-phase methylchlorosilanes. The split silicon peak was previously observed on Cu0 95i01 [21], Pd2Si [22,23] and Au—Si alloy [24,25]. Split Auger peaks were also reported for Al, P and S alloyed with copper [21]. Similarly, we observed a split peak (120 and 124 eV) for the phosphorus impurity in Cu0 95Si005. Hiraki et al. [21] suggested that the split peak is due to restructuring of the 3s—3p band of the third-period element because of hybridization with the 3d band of copper. Thus, according to Roth [231and Cros et al. [24,25], the silicon atoms in the alloy are in a more-metallic bonding environment which results in weaker silicon bonds.
370
T. C. Frank, fL. Falconer
/
Surface compositions of Cu
—
Si alloys
The weakened silicon bonding, observed by Cros et al. [24,25] when gold was evaporated onto Si(l 11), enhanced silicon oxidation at room temperature to form silicon dioxide [24,25]. Similarly, Cu3Si oxidized much more rapidly than silicon. Since Cu3Si also oxidized more rapidly than Cu0 95Si0 ~ the silicon bonds may be weaker in Cu3Si than in Cu095Si005. Indeed, in the direct-synthesis reaction, the reaction rate has been correlated with the amount of Cu3Si present [1,2]. Also, Gruhl and Engelmann [26] found that Cu3Si oxidized at 330—570 K to form Si02 faster than alloys with higher silicon concentrations. However, they [26] detected no visible corrosion for Cu3Si; ion bombardment may have removed a passivating layer in our samples to further increase the oxidation rate. 4.2. The Cu—Si surface
The surface composition estimates in table 3 are the same for the Cu(9l7 eV) peak (1.5 nm escape depth [3]) and the more surface-sensitive copper peaks at 60 and 105 eV (0.5 nm escape depth [3]). Thus, the top 1.5 nm of Cu—Si must be fairly uniform. Annealing Cu3Si above 770 K caused additional silicon to segregate to the surface and the more surface-sensitive copper peaks indicated higher silicon concentrations than did the Cu(9 17 eV) peak. Thus, the topmost layer in Cu3Si is significantly enriched in silicon after annealing. Some of the silicon is in an unalloyed form after annealing since the Si Auger peak was not split but was present as a broad peak at 92 eV. Thus, both alloyed and unalloyed silicon are present on the surface. The Cu—Si alloy has some similarities to Au—Si alloys. Both are noble metal—silicon systems with split silicon peaks and increased oxidation rates over elemental silicon. Cros et al. [24,25] also observed that annealing Au—Si above 670 K changed the split Si peak to a single peak at 92 eV. Thus, unalloyed silicon formed following this annealing. However, Cros et al. [24,25] observed that following annealing gold crystallites formed on top of a Au—Si monolayer and unalloyed silicon accumulated below. This ms in contrast to our results for Cu3Si. Annealing Cu0 95Si0 05 above 670 K caused no change in the split silicon peak. Also, Cu0 95Si0 ~ was not oxidized as rapidly as Cu3Si. Thus, significant differences exist between the surfaces of ordered-compound Cu3Si and the dilute alloy. 4.3. Surface composition
The Cu0 955i0 ~ alloy was studied to compare the surface composition of a dilute, solid-solution alloy with that of the ordered compound, Cu3Si. Silicon segregates to the surface of both alloys at 570 K. However, silicon segregation increased above 670 K for Cu3Si, apparently by forming both unalloyed and
T. C. Frank, fL. Falconer
/
Surface compositions of Cu
—
Si alloys
371
alloyed silicon. Annealing above 670 K did not change the Cu3Si bulk since ion bombardment always yielded a surface with the same bulk composition, with the split Si(90, 94 eV) Auger peak. Annealing Cu095Si005 at 770 K did not change the surface composition significantly. 4.3.1. Surface composition theories for solid-solution alloys
The predictions of several models of surface segregation will be compared to the experimental observation that silicon segregates to the surface of Cu095Si005. Bond-breaking theories 127,28] predict copper enrichment at the surface of Cu—Si alloys since copper has the lower heat of sublimation (330 versus 440 kJ/mol) [29], and the larger number of nearest neighbors (12 versus 4). Elastic-strain models predict no surface enrichment of solute atoms if solute atoms are smaller than solvent atoms [30]. Since the atomic radius of silicon is 10% smaller than that of copper [31], silicon segregation is not predicted. However, other theories predict that for dilute, ideal solutions for which each component has the same molar surface area, the component with the lower solid surface tension is enriched at the surface [32,33]. Silicon has a lower solid surface tension than copper (1.24 versus 1.67 N/rn) [32,34]. Thus, for Cu—Si solid-solutions, surface enrichment is predicted by the difference between surface tensions of pure components and not by bondstrength differences. Abra~iamand Brundle [35] pointed out that surface tension accounts for two effects: atoms occupying surface positions should have (1) a low bond energy and (2) a large atomic surface area. In other words, at equilibrium the surface tension y dG/3A is minimized. The component with higher bond energy can be enriched at the surface provided it has the larger atomic surface area. Silicon appears to have the higher surface area. The face-centered cubic structure of copper2/atom. has surface areas for thereconstruction, three low-inSi(1 11), without dex faces of 5.7, 6.6 and 9.3 X l0~20 m has a higher surface area, 12.8 X 1020 m2/atom [36]. The models discussed above are based on pure component properties. For better estimates of surface enrichment, alloy bonding and structure must be known. Burton and Machlin [37] used bulk phase diagrams to predict that if solute is enriched in the liquid phase, then solute will segregate to the surface provided that the separation between the solid and liquid phase lines is sufficiently large. Tsai et al.. [38] introduced the coefficient k CL/CS where CL is the solute concentration in the liquid phase that is in equilibrium with a 5% solute concentration in the solid phase (C 5 0.05). Solute surface enrichment is expected for K ~ 1.6 [38]. For the Cu—Si system, k = 2.0, indicating that silicon should segregate to the surface of dilute Cu—Si alloys, as was observed for our Cu0 95Si0)5 sample. The models discussed above predict that surface composition will approach the bulk composition as temperature increases. No change in surface composi=
=
=
372
T. C. Frank, J.L. Falconer
/
Surface compositions of Cu — Si alloys
tion was observed for Cu0 955i005 between 570 and 770 K. However, our ability to observe small changes in silicon enrichment is limited by experimental uncertainty of ±15% of the peak height ratios. Furthermore, the peak ratios in table 2 for Cu0 95Si005 are larger at 770 K than at 570 K for both copper and silicon spectra. This may be due to an increase in atomic density or the surface may be smoother at 770 K, since Auger signals are attenuated by surface roughness [3]. The estimated surface composition values shown in table 3 are identical only because the method used to estimate surface composition requires that atomic fractions sum to one. If the surface density is greater at 770 K than at 570 K, the F factor used to calculate composition is reduced and the calculated silicon atomic fraction is reduced. However, for a 20% increase in surface density, the composition only decreases one percentage point and the data are not accurate enough to measure this. 4.3.2. Surface composition theory for ordered alloys
The surface composition of ordered alloys is a function of crystal structure and orientation, temperature and the magnitude of a nearest-neighbor bonding parameter [39,40]. Van Santen and Sachtler [39] concluded that the surface of an ordered, binary alloy should be enriched in the component with the lower heat of sublimation provided that strain effects can be neglected and the quantity kT is less than twice the bonding parameter a E12 ~(E11 + E22). The symbol E,1 represents the bond energy between components i and j. The model was in agreement with their study of the Pt3Sn surface which was enriched in tin, the component with the lower heat of sublimation. A rough estimate of a is given by z~Hm/( X1 X2 z). The heat of mixing, ~ Hm, is not known for Cu3Si; however, comparison with similar alloys suggests that a rough estimate is 8 kJ/mol [10]. A value of 6 is used for z since the rhombohedral structure given by Solberg [9] for the Cu3Si phase is nearly cubic. This gives the value a 7 kJ/mol. The temperature at which kT 2a is 1700 K. Since the temperatures used in our study were much lower and since copper has a lower heat of sublimation than silicon, the model incorrectly predicts copper enrichment at the surface. Both Pt3Sn and Cu3Si are ordered alloys of a transition metal and a Group IVA element. In each case, the Group IVA element has a lower surface tension than the transition metal and the Group IVA element is enriched at the surface. It is possible that Cu3Si forms an amorphous surface layer on top of the bulk lattice since silicon enrichment is predicted by the solid-solution model. Morán-Lôpez and Bennemann [411 studied the ordering—disordering phasetransition at the surface of ordered fcc A3B-type alloys. For the fcc lattice, they showed that within the first few monolayers atomic ordering is reduced but not totally random. They predicted enrichment of component A in the first =
—
~‘
=
=
T C. Frank, f. L. Falconer
/
Surface compositions of Cu
—
Si alloys
373
atomic layer and enrichment to a smaller extent of component B in the second atomic layer.
5. Conclusions The surfaces of Cu3Si ordered compound and Cu095Si005 alloy were studied as a first step toward understanding the copper-catalyzed reaction between methyl chloride and silicon. The results for two separate Cu3Si samples are in good agreement. These studies show: The surfaces of both Cu3Si and Cu0 95Si005 are enriched in silicon, which is present in the alloyed form at 570 K. Above 670 K, additional unalloyed elemental silicon is present on the Cu3Si surface. The Cu3Si surface may be amorphous. Surface tension differences and the Burton and Machlin rule both predict silicon segregation to the surface of dilute, solid-solution Cu—Si alloys. Copper significantly alters silicon bonding upon alloying, as evidenced by the split Si(92 eV) Auger peak. The split peak makes identification of alloyed silicon possible. Clean Cu3Si oxidizes much more rapidly in air than does pure silicon due to weaker silicon bonds in the alloy. This weaker bonding may explain copper’s ability to catalyze the reaction between methyl chloride and silicon to form methylchlorosilanes. —
—
—
—
Acknowledgements We gratefully acknowledge support of this work under National Science Foundation Grant CPE8O-24236. We also thank the National Science Foundation for equipment grants ENG76-14388 and CPE79-23208 and the College of Engineering for equipment grants. We thank Union Carbide Corporation, Silicones and Urethane Intermediate Division, for partial support of this work, and the General Electric Company, Silicone Products Division, for support in the earlier stages of this work. We are very grateful to Drs. Jim C. Mikkelson and Kenrick M. Lewis for donation of alloy samples and Dr. Jim C. Mikkelson for donation of the sample preparation oven. We also thank Professor Seymour Geller for the X-ray analysis, Willy Grothe for construction of the vacuum system, and Jean Yves Levier and Joan M. Faiks for initial experimental design and improvements. We also appreciate the valuable discussions with Drs. Keith B. Kester and Kenrick M. Lewis.
374
T C. Frank, J.L. Falconer
/
Surface compositions of Cu
—
Si alloys
References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10)
[11] [12) [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37] [38] [39] [40] [41]
R.J.H. Voorhoeve, Organosilanes: Precursors to Silicones (Elsevier, Amsterdam. 1967). Al. Gorbunov, A.P. Belyi and G.G. Filippov, Russ. Chem. Rev. 43 (1974) 291. A.W. Czanderna, Ed., Methods of Surface Analysis (Elsevier, Amsterdam, 1975) ch. 5. T.C. Frank, MS Thesis, University of Colorado (1981). J.C. Mikkelson, Jr., private communication, 1980. B.H. Kolster, Acta Cryst. 19 (1965) 1049. V.D. Krylov, V.1. Agafonov and 0.1. Cherenkova, Soviet Phys.-Cryst. 11(1967)699. K.P. Mukherjee, J. Bandyopadhyaya and K.P. Gupta, AIME Trans. 245 (1969) 2335. J.K. Solberg, Acta Cryst. A34 (1978) 684. R. Hultgren, D.D. Pramod, D.T. Hawkins, M. Gleiser and K.K. Kelley, Selected Values of the Thermodynamic Properties of Binary Alloys (American Society for metals, Metals Park, OH, 1973). B.H. Kolster, J.C. Vlugter and R.J.H. Voorhoeve, Rec. Tray. Chim. 83 (1964) 737. L.E. Davis, N.C. MacDonald, P.W. Palmberg, G.E. Riach and RE. Weber, Handbook of Auger Electron Spectroscopy, 2nd ed. (Physical Electronics, Eden Prairie. MN, 1976). S.S. Chao, E.A. Knabbe and R.W. Vook, Surface Sci. 100 (1980) 581. P.M. Hall and J.M. Morabito, Surface Sci. 83 (1979) 391. Si. White, D.P. Woodruff and L. McDonnell, Surface Sci. 72(1978)77. M. Iwami, S.C. Kim, Y. Kataoka, T. Imura, A. Hiraki and F. Fujimoto, Japan. J. AppI. Phys. 19(1980)1627. A. Hiraki, S.C. Kim, T. Imura and M. Iwami, Japan. J. AppI. Phys. 18 (1979) 1767. J.G. McCarty and R.J. Madix, Surface Sci. 54 (1976) 121; J.G. McCarty, PhD Dissertation, Stanford University (1974) 108. H.F. Dylla, J.G. King and M.J. Cardillo, Surface Sci. 74 (1978) 141. KY. Yu, D.T. Ling and WE. Spicer, J. Catalysis 44 (1976) 373. A. Hiraki, S. Kim, W. Kammura and M. Iwami, Appl. Phys. Letters 34(1979)1. G.V. Robinson, AppI. Phys. Letters 25 (1974) 158. J.A. Roth, in: Proc. Symp. on Thin Film Phenomena — Interfaces and Interactions, Vol. 78-2, Eds. J.E.E. Baglin and J.M. Poate (Electrochemical Soc., Princeton, NJ, 1977) p. 29. A. Cros, F. Salvan, M. Commandre and J. Derrien, Surface Sci. 103 (1981) L109. A. Cros, J. Derrien and F. Salvan, Surface Sci. 110 (1981) 471. W. Gruhl and S. Engelmann, Metall 12 (1958) 985. J.J. Burton and R.L. Garten, Eds., Advanced Materials in Catalysis (Academic Press, New York, 1977) p. 33. W.H.H. Sachtler and R.A. van Santen, in: Advances in Catalysis, Eds. D.D. Eley, H. Pines and P.B. Weisz (Academic Press, New York, 1977) p. 69. Y.S. Touloukiau, Thermophysical Properties of High Temperature Solid Materials, Vol. 1 (MacMillan, New York, 1967). D.F. Ollis, J. Catalysis 59 (1979) 430. C.J. Smithells, Ed., Metals Reference Book, 5th ed. (Butterworths, London, 1976) p. 100. GA. Somorjai, Chemistry in Two Dimensions (Cornell Univ. Press, Ithaca, NY, 1981) pp. 31. 100. S.H. Overbury, G.A. Bertrand and GA. Somorjai, Chem. Rev. 75 (1975) 547. A.W. Adamson, Physical Chemistry of Surfaces (Wiley, New York, 1976) p. 269. F.F. Abraham and C.R. Brundle, J. Vacuum Sci. Technol. 18 (1981) 506. J.F. Nicholas, An Atlas of Models of Crystal Surfaces (Gordon and Breach, New York, 1965) pp. 29, 151. ii. Burton and ES. MachIm, Phys. Rev. Letters 37 (1976) 1433. N.H. Tsai, G.M. Pound and F.F. Abraham, J. Catalysis 50 (1977) 200. R.A. van Santen and W.M.H. Sachtler, J. Catalysis 33 (1974) 202. J.L. Morán-López and L.M. Falicov, Phys. Rev. B18 (1978) 2542. J.L. Morkn-López and K.H. Bennemann, Phys. Rev. B15 (1977) 4769.
SURFACE COMPOSITIONS OF COPPER-SILICON ALLOYS Timothy C. FRANK and John L. FALCONER Department of Chemical Engineering Campus Box 424, University of Colorado, Boulder, Colorado 80309, USA Received 3 September 1982; accepted for publication 13 December 1982
Auger spectroscopy of copper—silicon alloys showed that silicon is enriched at the surface of both an ordered compound, Cu 3Si, and a solid-solution alloy (5 at% Si). The surface compositions are estimated using elemental standards. Silicon segregation is predicted by thermodynamic theories based on surface tensions of pure components and by the empirical Burton and MachIm rule. The Si(92 eV) Auger peak splits into two peaks upon alloying, making identification of alloyed silicon possible. Surface silicon is in an alloyed form for both Cu3Si and Cu0 95Si005. Above 670 K, additional silicon segregates to the surface of Cu3Si and both alloyed and elemental silicon are present. Alloying copper with silicon results in weakened silicon bonds and clean Cu3Si oxidizes much more rapidly in air than does elemental silicon. This weakened bonding may be one reason that copper catalyzes the reaction between methyl chloride and silicon to form methylchlorosilanes.
I. Introduction In the commercial production of silicone polymers, organohalosilanes are hydrolized to form silanols, which condense to form silicone polymers and water. The principal industrial synthesis of these silanes is the direct-synthesis reaction between gaseous organic halides and elemental silicon. The most important reaction in this class of solid-catalyzed, gas—solid reactions is the formation of dimethyldichlorosilane ((CH3 ) 2SiCI2) using a copper catalyst. Cu
2 CH3C1
+
Si
—*
(CH3)2SiCI2.
(1)
Several studies have demonstrated that catalytic activity for this reaction is correlated with the amount of ordered-compound Cu3Si (~phase) present [1,21,but the reaction is poorly understood. As a first step toward understanding the reaction mechanism, we measured the surface compositions of clean ~ alloys as a function of temperature using Auger spectroscopy. Two Cu3Si samples, prepared separately, were used to determine the accuracy of the Auger measurements. A dilute, solid-solution 0378-5963/83/0000—0000/$03.OO © 1983 North-Holland
360
T C. Frank, J. L. Falconer
/
Surface compositions of Cu
—
Si alloys
alloy (5 at% Si in Cu) was also compared to Cu3Si to determine any surface differences that might influence the direct-synthesis reaction. Auger spectroscopy should yield a good estimate of surface composition for copper—silicon alloys since both copper and silicon have Auger peaks in the energy region which corresponds to minimum escape depth [3]. The silicon Auger peak at 92 eV split into two peaks upon alloying. This split peak serves as a good indication of the presence of alloyed silicon. However, it also complicates surface analysis, and surface compositions were estimated using the high-energy silicon peak and the copper peaks. The enrichment of silicon on the surfaces of these alloys was compared to theory predictions.
2. Experimental Surface analysis was carried out in an ion-pumped UHV chamber equipped with a PHI CMA Auger spectrometer, a UTI quadrupole mass spectrometer, and a Varian ion bombardment gun. After bakeout, pressures of 7 x i0~Pa (5 x l0~b0Torr) were easily obtained. Samples were either mounted on a PHI rotatable sample manipulator or in a transfer mechanism which allowed samples to be changed without exposing the UHV chamber to atmospheric pressure [4]. Samples were heated from behind by radiation from a 0.25 mm diameter tungsten filament. The temperature was measured with a chromel—alumel thermocouple (0.13 mm diameter wire) spot-welded to the side of the sample. X-ray analysis of the Cu3Si alloy was done using a Buerger precession camera. A Zeiss polarized-light microscope was used to observe the optical anisotropy of the samples. Two Cu3Si samples were studied. Sample No. 1 was prepared from a melt of 23 at% Si and 77 at% Cu since silicon precipitates formed at higher silicon concentrations [5]. The alloy was made from large chunks of silicon (6 mm diameter) in order to reduce the amount of Si02 present in the melt. The copper and silicon were sealed in an evacuated quartz ampoule and heated to 1300 K in a horizontal tube furnace. Vigorous agitation of the melt/solid mixture for 20 mm was required to incorporate unreacted silicon into the melt [5]. The resulting ingot was annealed at 670 K. Sample No. 2 was obtained from Union Carbide. The Cu0 95Si005 sample was also prepared by heating copper and silicon in an evacuated quartz ampoule at 1300 K in a horizontal tube furnace. The melt was agitated for 20 mm and then cooled to room temperature in 30 mm. Samples were sliced with a diamond saw and polished with Linde B abrasive (0.05 ~m Al 203). The Cu3 Si alloys were extremely brittle, so fairly thick samples (1.5 mm) were used. The Cu095Si005 sample was not brittle2and on was eacheasily face. sliced. The surface area of each sample was approximately 1 cm
T. C. Frank, J.L. Falconer
/ Surface
compositions of Cu — Si alloys
361
The silicon sample was an n-type single crystal, Si( 100), manufactured by Monsanto and containing iO-~%phosphorus. The copper was a polycrystalline foil (99.995% purity) supplied by Alfa Products. Argon used for ion bombardment was specified as 99.9995% pure (Scientific Gas Products, research grade) with less than 1 ppm H20 and 02. Carbon monoxide (Scientific Gas Products, technical grade) used for adsorption was 99% pure. All Auger spectra were recorded with a 3 keV beam voltage and a 30 .tA beam current. The modulation voltage was 6 V peak-to-peak for Si( 1619 eV) and I V peak-to-peak for all other peaks. The analyzer was calibrated using the negative peak of the 2000 eV, differentiated elastic peak.
3. Results Both the bulk and the surface of the samples were analyzed. X-ray diffraction and polarized-light microscopy were used for the bulk analysis. Auger spectroscopy, CO adsorption and scanning electron microscopy were used for surface analysis. The Auger spectra were recorded as a function of annealing temperature. Auger spectra of elemental silicon and copper were used in estimating surface compositions. 3.1. Bulk characterization
The X-ray diffraction for a small crystallite taken from the polycrystalline Cu3 Si sample No. 1 indicated an ordered structure with six-fold symmetry and hexagonal lattice constants a 0.70 nm and c 1.46 nm. The complex structure of Cu3Si is uncertain [6—9].Three temperature-dependent phases of Cu3Si are known to exist [10]; however, the active form for the direct-synthesis reaction is not known. The structure of our sample may be the mid-temperature phase (700—830 K) given by Solberg [9]; phase This is rhombohedral tR this 95.72°. corresponds R3 to with latticelattice constants aR 0.472 nm and hexagonal constants a 0.700 nm and c 0.732 nm. The six-fold symmetry for our alloy does not agree with the three-fold symmetry of Solberg’s structure but this may be due to twinning of the crystallite, since Cu 3Si has a tendency for twin formation [6,9]. The Cu3Si stoichiometry was verified on polished samples by using a polarized-light microscope. The three i~phases and the K phase (10—15 at% silicon) are the only optically-anisotropic compounds in the copper—silicon system. The K phase is not present at the conditions used to prepare our alloys [11]. The index of refraction for an optically-anisotropic material is a function of crystal orientation. When the polished Cu3Si surfaces were rotated between crossed polarizers, the intensity changed periodically, alternating between =
=
~‘
=
=
=
=
362
T. C. Frank, J.L. Falconer
/
Surface compositions of Cu — Si alloys
extinction and full illumination of the surface crystallites. Intensity does not change periodically for isotropic materials. This verifies the Cu3Si stoichiometry of our alloys. The polarized-light microscope showed that the Cu0 95Si0 ~ alloy was optically isotropic.
Si,570K
~3Si,77OK
dN(E) dE
A
Cu3SI,570K
~,5,770
40
K
60 80 100 120 Electron Energy (eV)
Fig. 1. Low-energy Auger spectra of alloys and pure elements after annealing at indicated temperatures for 30 mm.
T C. Frank, J.L. Falconer / Surface compositions of Cu — Si alloys
363
3.2. Surface characterization 3.2.1. Scanning electron micrographs
Before placement in the UHV chamber, the samples were polished to a bright mirror finish and analyzed by SEM. The only features observed in SEM were a few defects at magnifications of 17000 x After removal from the UHV chamber following repeated ion bombardment/annealing cycles, a roughened surface was observed on the Cu3Si samples. Scanning electron micrographs .
showed pitted areas 0.5 ~smwide and several hundred ~.tmlong. These pits were aligned and comprised about 10% of the total surface. 3.2.2. Surface cleaning
Cycles of argon ion bombardment and annealing were used all the 2 into6.7clean X l0~ Pa samples.Annealing Bombardment was570 done 500principal eV and 2—5 ~sA/cm were carbon and argon. was at K. atThe impurities oxygen, but nitrogen and sulfur were present on Cu 3Si and chlorine and sulfur were seen on copper. Long cleaning times reduced all contaminants below the AES detection limits on Cu3Si and copper. Small carbon and oxygen levels were still present on Cu0 95Si0 05 after 15 h of cycling. The 0(508 eV)/Cu(60 eV) and C(273 eV)/Cu(60 eV) Auger peak ratios were 0.05 and 0.02 respectively. Annealing Cu0 95Si005 to 770 K caused small concentrations of phosphorus and sulfur to segregate to the surface. Similar small carbon, oxygen and argon ratios were observed for the silicon sample. 3.2.3. Auger spectra
The pure silicon spectrum was very similar to that reported in the literature [12]; peak locations were within 1 eV. Auger peak energies for pure copper were also similar to those reported in the literature except for small differences in peak locations [12,13]. Several changes in the Auger spectra occurred on alloying, however. Fig. 1 shows the low-energy Auger spectra of copper, silicon, and the alloys. The low-energy silicon peak was significantly altered by alloying. Pure silicon has a single peak at 92 eV while both copper—silicon alloys exhibited two peaks at 90 and 94 eV. The peak at 82 eV for pure silicon shifted to 85 eV after alloying and its relative peak height increased from 0.01 to 0.15. The relative peak height is the ratio of secondary to primary peak height in a given peak ensemble (see table 1). Relative peak heights for the other silicon peaks also changed on alloying. The 1601 eV peak shifted to 1596 eV and its relative peak height decreased from 0.32 to 0.10 (see table I). The copper peaks, however, were not altered by alloying. Auger spectra of Cu3 Si were run both before and after silicon spectra, so that accurate comparisons could be made. All alloy spectra were reproducible within ±5—10% of peak height and ±1 eV in peak location. The shift in Auger peak locations are not due to inaccurate calibra-
364
T. C. Frank, J.L. Falconer
/ Surface
compositions of Cu — Si alloys
Table I Relative silicon peak heights Elemental silicon
Cu
3Si alloy
Peak location (eV)
Relative peak height ~
Peak location (eV)
92~ 82 16l9~ 1601 1582 1561 1543 1520
I 0.01 I 0.32 0.13 0.27 0.12 0.05
90,94~> 85 l620~ 1596 1579 1560 1538 1518
Relative peak height a)
0.15 0.10 0.10 0.25 0.05 0.10
~ Relative peak height; the ratio of secondary peak height to primary peak height in a given peak ensemble b) The reference peak used to calculate the relative peak heights for energies that follow in the table. ~ Si(94 CV) is used as the reference peak for the alloy.
tion; the copper peaks at 60 and 917 eV were the same for both copper and the alloys. The two silicon peaks at 90 and 94 eV changed to a single peak at 92 eV when the Cu3 Si alloy was annealed above 670 K, as shown in fig. 1. This single Si(92 eV) peak is much broader than the corresponding peak measured on pure silicon and may be a combination of spectra for alloyed and unalloyed silicon. The split silicon peak measured on Cu0 95Si005 did not change to a single peak when this alloy was heated between 300 and 830 K. Fig. 2 shows peak height as a function of annealing temperature for Cu3Si alloy No. 2. The data were obtained by annealing the ion-bombarded sample at the temperatures indicated in fig. 2, beginning at 370 K and progressing to higher temperatures. Auger spectra were recorded after sufficient annealing times that the Auger peak heights reached steady-state values: 30 mm at 370 K, 20 mm at temperatures up to 670 K, and 10 mm at higher temperatures. In separate experiments, the sample was annealed for 145 mm at 370 K, 60 mm at 470 K, 60 mm at 570 K, and 30 mm at 770 K with the same results as in fig. 2, indicating that the equilibrium surface was rapidly attained. 3.2.4. Surface composition
The ratio of an Auger peak height in the alloy to the same transition peak height in the pure component should estimate the surface concentration of that component in the alloy. Table 2 lists these ratios. With few exceptions the
T C. Frank, J. L. Falconer
I
I
I
/
Surface compositions of Cu — Si alloys
I
I
Si
90,94
5
,,—~~ci—-—a-Si92
Cu6~
~°
-
SI9~
~
~
Si~6~0~
~
Cu917
-
Si1620
-
CX 15)
0
-(X15) 2
-
~
-
~
I
Si92
-
I~
0
-
~Cu60
-
-
-
Si90,94
-
1
-
split height —i
-
0
-
300
365
I 400
500 600 Temperature (K)
I 700
I 800
-
Fig. 2. Equilibrium Auger peak heights for Cu3Si sample as a function of annealing temperature.
agreement between values obtained on samples No. 1 and No. 2 of Cu 3Si is very good. Also, for most ratios, the silicon atomic fractions estimated from the copper peaks (by subtracting the copper peak ratio from one) are in reasonable agreement with the fractions estimated from Sm( 1619 eV). Following annealing of Cu3Si at 770 K, the more surface-sensitive copper Auger peaks indicate a much higher silicon concentration than the high-energy copper or silicon peaks. However, for Cu 0.95 Si 0.05’ annealing at 770 K increased all the Auger peaks, complicating estimates of surface composition. For all samples, the Si(94 eV) peak indicates a silicon surface concentration 1.5 to 2.5 times as large as that obtained from the other peaks. Since this silicon Auger peak shape changed upon alloying, comparisons using this peak are not meaningful. To obtain better estimates of surface compositions, the method of Hall and Morabito was used to account for matrix effects [14]. This method uses a correction factor F which accounts for differences between Auger transitions in the alloy and pure silicon. This correction is for dilute alloys, evaluated in the limit of pure copper, and is independent of the copper transition used. This method is not strictly valid for concentrated alloys such as Cu3Si, but F is not
366
T. C. Frank, J.L. Falconer
/
Surface compositions of Cu — Si alloys
aV 0.
V .-
C
9
N
~
V
V
V
~o o’~~——~ ~a N — ~ ddd~~~d~ .2 C
-~
V — > V
0
N
0.
>
N-N’C~0’N’C ~0
U
~O
‘~ 0’
N 0’
‘O
~
a
a -o
‘5
0
V
‘c~ C
.9 I-
Zr.. ~
U
N s ~
N ‘.0 ‘.0
~ ‘.0
~O N en 0’
~
~ 0’
a ~~
.~
—
I ~ 0.
~0
<
0.
v
C
U0
V
VO’ — N
~
0’ 0’ ‘t~ ‘0
— ‘i”.
~n
‘0 ~O
‘
C ~
.5 —
~
t
E
‘0 ‘0
~
C N
‘0
U
CC N N r’. N
‘0
‘0 ~ ‘0
5
CC N N
0,
—
~a .2
V
~
o~oaa0~~ a
~
cc
cc
.OQCUQ.0QV
~
~
a a
— ~0 —
.0
a 0’
0
U
Z
N
0
Z
..~
N
a
0
U
U
U
.0
T. C. Frank, J.L. Falconer
/
367
Surface compositions of Cu — Si alloys
a strong function of composition [14] and the atomic densities of Cu and Cu3Si 3) [7]. The method assumes are and and 7.8 Xthat 1022 atoms/cm peaksimilar shapes(8.5 do x not1022 change surface roughness and diffraction effects are negligible. Diffraction errors should be small since Auger emission from Si(100) is nearly isotropic [15] and other samples are polycrystalline. Hall and Morabito [14] used two correlations for escape depths in correcting for matrix effects. When corrected for the 3 keV primary beam voltage used in this study, the values for F are 0.665 (Penn correlation) and 0.905 (Seah and Dench correlation) [14]. Table 3 lists the silicon atom fractions obtained with the Hall and Morabito matrix corrections for both Penn’s and Seah and Dench’s (SD) correlations for escape depth. All calculations are based on the Si(l6l9 eV) transition since the more surface-sensitive transition at 92 eV changes shape with alloying. For the ion-bombarded samples, surface composition estimates using the Penn correlation are very close to bulk values. Since previous studies have indicated that sputtering yields of copper and silicon in Cu—Si alloys are equal [16,17], the ion-bombarded sample should have a surface composition close to the bulk. Thus the Penn correlation may yield the more accurate estimates. However, it is possible that the Auger beam, which heated the samples to 350—400 K, annealed the surface and caused silicon enrichment. Table 3 shows that following annealing at 570 K, the surface compositions of Cu 3Si and Cu0 955i005 are at least 30 at% Si and 10 at% Si, respectively. Agreement between the two Cu3 Si samples is very good. Annealing Cu3 Si at
Table 3 Silicon atomic fraction a) for various copper peaks Sample
Surface treatment
Cu3Si (No. 1) Cu3Si (No. 2)
Cu0 95Si005
a)
Ion bombarded Annealed at 570 Ion bombarded Annealed at 570 Annealed at 770 Ion bombarded Annealed at 570 Annealed at 770
K K K K K
Cu(60 eV)
Cu(105 eV)
Cu(917 eV)
Penn
SD
Penn
SD
Penn
SD
0.24 0.30 0.29 0.33 0.50 0.07 0.10 0.10
0.30 0.37 0.36 0.40 0.58 0.09 0.13 0.13
0.23 0.29 0.29 0.32 0.45 0.06 0.10 0.10
0.29 0.36 0.36 0.39 0.53 0.08 0.12 0.13
0.25 0.29 0.29 0.30 0.37 0.06 0.10 0.10
0.31 0.36 0.36 0.37 0.44 0.08 0.13 0.13
(paIIoY/pSi)F S~
(paIloY/pca)+
(p;11oY/p~)F
The Si(1619 eV) peak was used for all calculations.
368
T. C. Frank, J.L. Falconer / Surface compositions of Cu
—
Si alloys
770 K produced a surface containing more than 40 at% Si, with higher values indicated for the more surface-sensitive peaks. Annealing Cu0 95Si005 at 770 K caused no additional change in the surface composition though both copper and silicon peak heights increased (table 2). The silicon surface enrichment observed is not an artifact due to preferential sputtering during ion bombardment. Following bombardment, the surface composition was approximately the same as the bulk and enrichment only occurred on annealing. Also, additional bombardment yielded the same bulk composition. Moreover, no preferential sputtering was detected in previous studies of Cu—Si alloys [16,17]. 3.3. CO adsorption
Preliminary CO desorption experiments were carried out for Si, Cu and Cu3Si in an attempt to determine the composition of the surface monolayer. 2) for On three surfaces COK.coverage was lowadsorption ((1—2) x l0~molecules/cm 50 Lalladsorption at 310 No additional was observed at higher exposures. Carbon monoxide adsorption on nickel was used for calibration [18].On silicon and copper, a small CO peak was observed at 370 K. A similar
/1
/1
Contaminated Cu 3Si
dN(E)
Coppe~0PP&
273
Clean Cu3Si
~
Copp~~~pper
Carbon
60 67
67 79 85
Silicon
85 94
Silicon
508
90
Oxygen //
Electron Energy (eV) Fig. 3. Low-energy Auger spectra of contaminated and clean Cu3Si sample. The clean sample was annealed at 570 K.
T. C. Frank, fL. Falconer
/ Surface
compositions of Cu — Si alloys
369
peak has been observed on Si(100) [19].Carbon monoxide is weakely adsorbed on copper [20]. A broad peak appeared at 330 K on Cu3Si for samples annealed at 570 or 770 K. The CO desorption data could not be used to determine the surface monolayer composition because CO coverage was too low. The surface could have changed significantly without affecting CO desorption. 3.4. Cu3 Si oxidation
The ion-bombarded Cu3Si alloy oxidized much more readily than Cu0 95Si005, the pure components, or the Cu3Si samples before ion bombardment. The cleaned Cu3Si samples formed a black oxide surface layer after less than 24 h exposure to room-temperature air, while the pure components and the Cuo 95 0.05 alloy showed no visible corrosion even after several months exposure. Fig. 3 shows the Auger spectra of a clean Cu3Si alloy which was exposed to methyl chloride and oxygen and then heated. Oxygen and carbon were the only contaminants observed. The intensity of the peak at 67, 79 and 85 eV increased for the contaminated surface. This plus the oxygen peaks at 475, 490 and 508 eV may indicate the presence of Si02. The relative amplitudes of the 90 and 94 eV peaks also changed but both peaks were present, indicating surface silicon was still alloyed with copper.
4. Discussion 4.1. Silicon bonding in Cu—Si alloys
Copper significantly alters silicon bonding upon alloying, as evidenced by the split Si(92 eV) Auger peak, the changes in other silicon peaks, and the increased oxidation rate of Cu3Si. Silicon bonding appears to be weakened in the alloy, particularly in Cu3Si, and this may be one reason why copper, in Cu3 Si, is a good catalyst for the direct-synthesis reaction. The weakened silicon bond in the alloys is more easily broken to form gas-phase methylchlorosilanes. The split silicon peak was previously observed on Cu0 95i01 [21], Pd2Si [22,23] and Au—Si alloy [24,25]. Split Auger peaks were also reported for Al, P and S alloyed with copper [21]. Similarly, we observed a split peak (120 and 124 eV) for the phosphorus impurity in Cu0 95Si005. Hiraki et al. [21] suggested that the split peak is due to restructuring of the 3s—3p band of the third-period element because of hybridization with the 3d band of copper. Thus, according to Roth [231and Cros et al. [24,25], the silicon atoms in the alloy are in a more-metallic bonding environment which results in weaker silicon bonds.
370
T. C. Frank, fL. Falconer
/
Surface compositions of Cu
—
Si alloys
The weakened silicon bonding, observed by Cros et al. [24,25] when gold was evaporated onto Si(l 11), enhanced silicon oxidation at room temperature to form silicon dioxide [24,25]. Similarly, Cu3Si oxidized much more rapidly than silicon. Since Cu3Si also oxidized more rapidly than Cu0 95Si0 ~ the silicon bonds may be weaker in Cu3Si than in Cu095Si005. Indeed, in the direct-synthesis reaction, the reaction rate has been correlated with the amount of Cu3Si present [1,2]. Also, Gruhl and Engelmann [26] found that Cu3Si oxidized at 330—570 K to form Si02 faster than alloys with higher silicon concentrations. However, they [26] detected no visible corrosion for Cu3Si; ion bombardment may have removed a passivating layer in our samples to further increase the oxidation rate. 4.2. The Cu—Si surface
The surface composition estimates in table 3 are the same for the Cu(9l7 eV) peak (1.5 nm escape depth [3]) and the more surface-sensitive copper peaks at 60 and 105 eV (0.5 nm escape depth [3]). Thus, the top 1.5 nm of Cu—Si must be fairly uniform. Annealing Cu3Si above 770 K caused additional silicon to segregate to the surface and the more surface-sensitive copper peaks indicated higher silicon concentrations than did the Cu(9 17 eV) peak. Thus, the topmost layer in Cu3Si is significantly enriched in silicon after annealing. Some of the silicon is in an unalloyed form after annealing since the Si Auger peak was not split but was present as a broad peak at 92 eV. Thus, both alloyed and unalloyed silicon are present on the surface. The Cu—Si alloy has some similarities to Au—Si alloys. Both are noble metal—silicon systems with split silicon peaks and increased oxidation rates over elemental silicon. Cros et al. [24,25] also observed that annealing Au—Si above 670 K changed the split Si peak to a single peak at 92 eV. Thus, unalloyed silicon formed following this annealing. However, Cros et al. [24,25] observed that following annealing gold crystallites formed on top of a Au—Si monolayer and unalloyed silicon accumulated below. This ms in contrast to our results for Cu3Si. Annealing Cu0 95Si0 05 above 670 K caused no change in the split silicon peak. Also, Cu0 95Si0 ~ was not oxidized as rapidly as Cu3Si. Thus, significant differences exist between the surfaces of ordered-compound Cu3Si and the dilute alloy. 4.3. Surface composition
The Cu0 955i0 ~ alloy was studied to compare the surface composition of a dilute, solid-solution alloy with that of the ordered compound, Cu3Si. Silicon segregates to the surface of both alloys at 570 K. However, silicon segregation increased above 670 K for Cu3Si, apparently by forming both unalloyed and
T. C. Frank, fL. Falconer
/
Surface compositions of Cu
—
Si alloys
371
alloyed silicon. Annealing above 670 K did not change the Cu3Si bulk since ion bombardment always yielded a surface with the same bulk composition, with the split Si(90, 94 eV) Auger peak. Annealing Cu095Si005 at 770 K did not change the surface composition significantly. 4.3.1. Surface composition theories for solid-solution alloys
The predictions of several models of surface segregation will be compared to the experimental observation that silicon segregates to the surface of Cu095Si005. Bond-breaking theories 127,28] predict copper enrichment at the surface of Cu—Si alloys since copper has the lower heat of sublimation (330 versus 440 kJ/mol) [29], and the larger number of nearest neighbors (12 versus 4). Elastic-strain models predict no surface enrichment of solute atoms if solute atoms are smaller than solvent atoms [30]. Since the atomic radius of silicon is 10% smaller than that of copper [31], silicon segregation is not predicted. However, other theories predict that for dilute, ideal solutions for which each component has the same molar surface area, the component with the lower solid surface tension is enriched at the surface [32,33]. Silicon has a lower solid surface tension than copper (1.24 versus 1.67 N/rn) [32,34]. Thus, for Cu—Si solid-solutions, surface enrichment is predicted by the difference between surface tensions of pure components and not by bondstrength differences. Abra~iamand Brundle [35] pointed out that surface tension accounts for two effects: atoms occupying surface positions should have (1) a low bond energy and (2) a large atomic surface area. In other words, at equilibrium the surface tension y dG/3A is minimized. The component with higher bond energy can be enriched at the surface provided it has the larger atomic surface area. Silicon appears to have the higher surface area. The face-centered cubic structure of copper2/atom. has surface areas for thereconstruction, three low-inSi(1 11), without dex faces of 5.7, 6.6 and 9.3 X l0~20 m has a higher surface area, 12.8 X 1020 m2/atom [36]. The models discussed above are based on pure component properties. For better estimates of surface enrichment, alloy bonding and structure must be known. Burton and Machlin [37] used bulk phase diagrams to predict that if solute is enriched in the liquid phase, then solute will segregate to the surface provided that the separation between the solid and liquid phase lines is sufficiently large. Tsai et al.. [38] introduced the coefficient k CL/CS where CL is the solute concentration in the liquid phase that is in equilibrium with a 5% solute concentration in the solid phase (C 5 0.05). Solute surface enrichment is expected for K ~ 1.6 [38]. For the Cu—Si system, k = 2.0, indicating that silicon should segregate to the surface of dilute Cu—Si alloys, as was observed for our Cu0 95Si0)5 sample. The models discussed above predict that surface composition will approach the bulk composition as temperature increases. No change in surface composi=
=
=
372
T. C. Frank, J.L. Falconer
/
Surface compositions of Cu — Si alloys
tion was observed for Cu0 955i005 between 570 and 770 K. However, our ability to observe small changes in silicon enrichment is limited by experimental uncertainty of ±15% of the peak height ratios. Furthermore, the peak ratios in table 2 for Cu0 95Si005 are larger at 770 K than at 570 K for both copper and silicon spectra. This may be due to an increase in atomic density or the surface may be smoother at 770 K, since Auger signals are attenuated by surface roughness [3]. The estimated surface composition values shown in table 3 are identical only because the method used to estimate surface composition requires that atomic fractions sum to one. If the surface density is greater at 770 K than at 570 K, the F factor used to calculate composition is reduced and the calculated silicon atomic fraction is reduced. However, for a 20% increase in surface density, the composition only decreases one percentage point and the data are not accurate enough to measure this. 4.3.2. Surface composition theory for ordered alloys
The surface composition of ordered alloys is a function of crystal structure and orientation, temperature and the magnitude of a nearest-neighbor bonding parameter [39,40]. Van Santen and Sachtler [39] concluded that the surface of an ordered, binary alloy should be enriched in the component with the lower heat of sublimation provided that strain effects can be neglected and the quantity kT is less than twice the bonding parameter a E12 ~(E11 + E22). The symbol E,1 represents the bond energy between components i and j. The model was in agreement with their study of the Pt3Sn surface which was enriched in tin, the component with the lower heat of sublimation. A rough estimate of a is given by z~Hm/( X1 X2 z). The heat of mixing, ~ Hm, is not known for Cu3Si; however, comparison with similar alloys suggests that a rough estimate is 8 kJ/mol [10]. A value of 6 is used for z since the rhombohedral structure given by Solberg [9] for the Cu3Si phase is nearly cubic. This gives the value a 7 kJ/mol. The temperature at which kT 2a is 1700 K. Since the temperatures used in our study were much lower and since copper has a lower heat of sublimation than silicon, the model incorrectly predicts copper enrichment at the surface. Both Pt3Sn and Cu3Si are ordered alloys of a transition metal and a Group IVA element. In each case, the Group IVA element has a lower surface tension than the transition metal and the Group IVA element is enriched at the surface. It is possible that Cu3Si forms an amorphous surface layer on top of the bulk lattice since silicon enrichment is predicted by the solid-solution model. Morán-Lôpez and Bennemann [411 studied the ordering—disordering phasetransition at the surface of ordered fcc A3B-type alloys. For the fcc lattice, they showed that within the first few monolayers atomic ordering is reduced but not totally random. They predicted enrichment of component A in the first =
—
~‘
=
=
T C. Frank, f. L. Falconer
/
Surface compositions of Cu
—
Si alloys
373
atomic layer and enrichment to a smaller extent of component B in the second atomic layer.
5. Conclusions The surfaces of Cu3Si ordered compound and Cu095Si005 alloy were studied as a first step toward understanding the copper-catalyzed reaction between methyl chloride and silicon. The results for two separate Cu3Si samples are in good agreement. These studies show: The surfaces of both Cu3Si and Cu0 95Si005 are enriched in silicon, which is present in the alloyed form at 570 K. Above 670 K, additional unalloyed elemental silicon is present on the Cu3Si surface. The Cu3Si surface may be amorphous. Surface tension differences and the Burton and Machlin rule both predict silicon segregation to the surface of dilute, solid-solution Cu—Si alloys. Copper significantly alters silicon bonding upon alloying, as evidenced by the split Si(92 eV) Auger peak. The split peak makes identification of alloyed silicon possible. Clean Cu3Si oxidizes much more rapidly in air than does pure silicon due to weaker silicon bonds in the alloy. This weaker bonding may explain copper’s ability to catalyze the reaction between methyl chloride and silicon to form methylchlorosilanes. —
—
—
—
Acknowledgements We gratefully acknowledge support of this work under National Science Foundation Grant CPE8O-24236. We also thank the National Science Foundation for equipment grants ENG76-14388 and CPE79-23208 and the College of Engineering for equipment grants. We thank Union Carbide Corporation, Silicones and Urethane Intermediate Division, for partial support of this work, and the General Electric Company, Silicone Products Division, for support in the earlier stages of this work. We are very grateful to Drs. Jim C. Mikkelson and Kenrick M. Lewis for donation of alloy samples and Dr. Jim C. Mikkelson for donation of the sample preparation oven. We also thank Professor Seymour Geller for the X-ray analysis, Willy Grothe for construction of the vacuum system, and Jean Yves Levier and Joan M. Faiks for initial experimental design and improvements. We also appreciate the valuable discussions with Drs. Keith B. Kester and Kenrick M. Lewis.
374
T C. Frank, J.L. Falconer
/
Surface compositions of Cu
—
Si alloys
References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10)
[11] [12) [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37] [38] [39] [40] [41]
R.J.H. Voorhoeve, Organosilanes: Precursors to Silicones (Elsevier, Amsterdam. 1967). Al. Gorbunov, A.P. Belyi and G.G. Filippov, Russ. Chem. Rev. 43 (1974) 291. A.W. Czanderna, Ed., Methods of Surface Analysis (Elsevier, Amsterdam, 1975) ch. 5. T.C. Frank, MS Thesis, University of Colorado (1981). J.C. Mikkelson, Jr., private communication, 1980. B.H. Kolster, Acta Cryst. 19 (1965) 1049. V.D. Krylov, V.1. Agafonov and 0.1. Cherenkova, Soviet Phys.-Cryst. 11(1967)699. K.P. Mukherjee, J. Bandyopadhyaya and K.P. Gupta, AIME Trans. 245 (1969) 2335. J.K. Solberg, Acta Cryst. A34 (1978) 684. R. Hultgren, D.D. Pramod, D.T. Hawkins, M. Gleiser and K.K. Kelley, Selected Values of the Thermodynamic Properties of Binary Alloys (American Society for metals, Metals Park, OH, 1973). B.H. Kolster, J.C. Vlugter and R.J.H. Voorhoeve, Rec. Tray. Chim. 83 (1964) 737. L.E. Davis, N.C. MacDonald, P.W. Palmberg, G.E. Riach and RE. Weber, Handbook of Auger Electron Spectroscopy, 2nd ed. (Physical Electronics, Eden Prairie. MN, 1976). S.S. Chao, E.A. Knabbe and R.W. Vook, Surface Sci. 100 (1980) 581. P.M. Hall and J.M. Morabito, Surface Sci. 83 (1979) 391. Si. White, D.P. Woodruff and L. McDonnell, Surface Sci. 72(1978)77. M. Iwami, S.C. Kim, Y. Kataoka, T. Imura, A. Hiraki and F. Fujimoto, Japan. J. AppI. Phys. 19(1980)1627. A. Hiraki, S.C. Kim, T. Imura and M. Iwami, Japan. J. AppI. Phys. 18 (1979) 1767. J.G. McCarty and R.J. Madix, Surface Sci. 54 (1976) 121; J.G. McCarty, PhD Dissertation, Stanford University (1974) 108. H.F. Dylla, J.G. King and M.J. Cardillo, Surface Sci. 74 (1978) 141. KY. Yu, D.T. Ling and WE. Spicer, J. Catalysis 44 (1976) 373. A. Hiraki, S. Kim, W. Kammura and M. Iwami, Appl. Phys. Letters 34(1979)1. G.V. Robinson, AppI. Phys. Letters 25 (1974) 158. J.A. Roth, in: Proc. Symp. on Thin Film Phenomena — Interfaces and Interactions, Vol. 78-2, Eds. J.E.E. Baglin and J.M. Poate (Electrochemical Soc., Princeton, NJ, 1977) p. 29. A. Cros, F. Salvan, M. Commandre and J. Derrien, Surface Sci. 103 (1981) L109. A. Cros, J. Derrien and F. Salvan, Surface Sci. 110 (1981) 471. W. Gruhl and S. Engelmann, Metall 12 (1958) 985. J.J. Burton and R.L. Garten, Eds., Advanced Materials in Catalysis (Academic Press, New York, 1977) p. 33. W.H.H. Sachtler and R.A. van Santen, in: Advances in Catalysis, Eds. D.D. Eley, H. Pines and P.B. Weisz (Academic Press, New York, 1977) p. 69. Y.S. Touloukiau, Thermophysical Properties of High Temperature Solid Materials, Vol. 1 (MacMillan, New York, 1967). D.F. Ollis, J. Catalysis 59 (1979) 430. C.J. Smithells, Ed., Metals Reference Book, 5th ed. (Butterworths, London, 1976) p. 100. GA. Somorjai, Chemistry in Two Dimensions (Cornell Univ. Press, Ithaca, NY, 1981) pp. 31. 100. S.H. Overbury, G.A. Bertrand and GA. Somorjai, Chem. Rev. 75 (1975) 547. A.W. Adamson, Physical Chemistry of Surfaces (Wiley, New York, 1976) p. 269. F.F. Abraham and C.R. Brundle, J. Vacuum Sci. Technol. 18 (1981) 506. J.F. Nicholas, An Atlas of Models of Crystal Surfaces (Gordon and Breach, New York, 1965) pp. 29, 151. ii. Burton and ES. MachIm, Phys. Rev. Letters 37 (1976) 1433. N.H. Tsai, G.M. Pound and F.F. Abraham, J. Catalysis 50 (1977) 200. R.A. van Santen and W.M.H. Sachtler, J. Catalysis 33 (1974) 202. J.L. Morán-López and L.M. Falicov, Phys. Rev. B18 (1978) 2542. J.L. Morkn-López and K.H. Bennemann, Phys. Rev. B15 (1977) 4769.