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Bilayers with chromium disilicide: Chromium-vanadium
~C6
IS~LAYERS WITH CHROMHJM
Applied Surface Science38 (1989) 106-116 North-Holland, Amsterdam
DISHL~CIIDE: CHROM~UM-VANADIIUM
T.L FINSTAD ~), O. THOMAS b.c) and F.M. D ' H E U R L E b.d) ~ Fystk Institutt, Universitetet i Oslo, Boks 1048, Blindern, 0316 Oslo 3, Norway bj I B M T.J. Watson Research Center. P O 218, Yorktown Heights, N Y 10598, USA ~) Laboratoire des Mat~riaux et de G~nie Physique, E N S P G , BP 46, 38402 Saint Martin d'Hdres, France •o Fasta Tillstdmdets Elektronik, K T H - Electrum, Box 1298, 164 28 Kista, Sweden
Received 20 March 1989; accepted for publication 27 April 1989
The behavior of chromium-vanadiumlayers deposited in that order on siliconsubstrates has been studied, as a function of heat treatment, with the aim of determining the difference of mobilitiesbetweenthe silicon and the metal atom.~in the two respectivedisilicides.The motion of the silicon atoms requires temperatures of the order of 500°C, that of the metal atoms temperatures about 300 °C higher.The two disilicides,previouslyreported to be totally soluble at 1500°C, appear to remain totally soluble at least down to 900°C.
I. lIntreduetion The thermal oxidation (meaning here the formation of a layer of pure SiO 2) of silicides is interesting [1-5] not only in itself, but also because it may throw light upon the still unresolved complexities of the oxidation of silicon. One aspect of this problem is the injection [6-8] of defects in the silicon during oxidation. Intriguing observations made on the oxidation of silicides seem to bear on this question. Duriog the oxidation of CoSi 2 and NiSi 2 one observes not only the decomposition [3,5] of the disilicides accompanied by the reverse motion (from the silicide-oxide interface to the silicon substrate) of the metal atoms, as during ;he formation of the disilicides, but also the motion in the same direction of about 25% of the silicon atoms which could have been oxidized. The contrast [3,4] between the formation of CrSi 2, whzn the silicon atoms constitute the most mobile species, and the oxidation of the same silicide, which occurs via the reverse motion of the chromium atoms, is equally surprising. The present experiments were undertaken in order to investigate the difference between the mobilities of the metal and the silicon atoms in CrSi 2. The experiments which have been mentioned [3,4] above were conducted with inert markers [9,10] according to an old and well established technique. This provides unambiguous but not so accurate results since they prove that CrSi 2 0169-4332/89/$03.50 © Elsevier Science Publishers B.V. (North-Holland Physics Publishing Division)
T.J. Finstad et al. / Bilayers with chrom m disihcide
107
forms predominantly by silicon motion, but ~ 'e may only say that the mobility of the metal atoms is less than about 1( + that of the silicon atoms. However, the ratio could be much smaller. In vie,,, of the suprising behaviors reported above, it becomes important to evaluate :ow big the difference might be. Namely, one wants to know as quantitatively ~ possible how different are the behaviors of the moving species during silicice formation and oxidation. To this end one proposes to study the rate at whicL similar silicides, here CrSi z and VSi 2 (one may also study pairs composed of CrSi 2 and MoSi z or W S i , ) in intimate contact with one another intermix. Th~s technique has been used to investigate the formations [11,12] of MoSi 2 and WSi z, as well as NbSi z and TaSi z. A detailed discussion is found in ref. [23]. For reasons which will become apparent later on the experiments reported here with vanadium-chromium bilayers are quite similar to experiments [14] conducted with tantalum-tungsten bilayers. In a slightly different spirit one may look also at oxidation studies carried out with niobium-tantalum bilayers [15].
2. Experimental procedures and results In order to minimize the formation of any interface oxide the two layers of chromium and vanadium were deposited in tiffs order in a single pump-down, in a vacuum chamber equipped with dual e-beam heated sources. The pressure during deposition was 1 × 10 -7 Toil', the substrate temperature 150°C: the rate of deposition 10 A / s , and the nominal film thicknesses 90 nm each (corresponding to respective disilicide thicknesses of 270 and 243 rim). The samples about 1 cm square cut from the 6 cm wafers were heat-treated in an atmosphere of helium purified over a titanium bed at about 8 5 0 ° C and zirconium at 650°C. Fig. 1, which displays several backscattering spectra, shows that the formation of CrSi2 is complete after annealing at 450 ° C for one hour, however, the formation of VSi2 requires a temperature of 600 o C. This is quite a good agreement with previously reported [16-18] results. The sensitivity of backscattering is limited for distinguishing elements with close atomic numbers such as chromium and vanadium. Thus in order to follow the effects of annealing at high temperatures one has to have recourse to Auger spectroscopy combined with sputter profiling, as was done with tantalumtungsten [14], and cobalt-nickel [19]. Supplemental information was obtained, when necessary, from X-ray diffraction patterns from a computer-controlled diffractometer equipped with a copper tube and a post-sample monochromator (which eliminates fluorescence radiations). Separate Auger spectra were obtained from CrSi 2 and VSi2, which were used as calibration standards. Three lines [20] were used for estimating the concentrations: C r L M M at about 524 eV, V L M M at about 469 eV, and Si KLL at about 1608 eV. Two of these are shawn in fig. 2, both for the
108
T.J. Finstad et aL
8/--
'
/
Bilayers with chromium disilicide
I
~
' Cr
A
V
I!0 f.4 eACKSCATTERING ENERGY (MoV)
1.8
Fig. l. Backscattering spectra of the chromlum-vanadium bilayers, respectively as-deposited and after heat treatments for one hour at 450 o C and 600 o C. Because the atomic number of vanadium is smaller than that of chromium (by one) the spectral part of vanadium is less high than that of chromium (both in the initial metallic slate and in the disilicide, after annealing at 600°C). For the same reason, and because the vanadium layer is found at the free surface the overlap of the chromium and vanadium counts causes a peak in the metal part of the spectra (at about 1.6 MeV initially, and at 1.5 MeV after completion of the disilicide reaction). Since the vanadium layer, both in the as-deposited state and immediately after silicidation, is found at the free surface the position of "surface vanadium" corresponds to the high energy side of the displayed spectra slightly below 1.7 MeV. Surface chromium would be found at a somewhat higher energy, with a vanadium-chromium separation of 11 keV for 2.3 MeV alpha particles; this would be hardly visible on the energy scale used here.
'
I
. . . . ,,,.%.~Z--
I
' ' ' " .~--'-'-Cr
,
I
' ' Si 2
I
.
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1
a
i
VANADIUM
,
,
I
400
.
.
.
.
t
,
45O
.,
,
.
50O
.
.
.
.
550
ENERGY (eV) Fig. 2. Partial Auger spectra obtained from standard VSi 2 and CrSi 2 samples and from the middle of the vanadium-rich and chromium-rich section of a bilayer sample annealed at 700 ° C for one hour. The small peak showing at low energy for one sample only is due to nitrogen contamination during the Auger analysis. The shaded areas were used to calculate alloy concentrations.
109
T.J. Finstad et aL / Bilayers with chromium disilicide
soo 7oo 8oo 9oo Iooo
o.4
c c c c c
oc~ oc~ ~Tcr oct &cr
@v ov ~v mv Av
o.,~ o.t o
'
DEPTH ($FUTTERI~ TI~E)
Fig. 3. Concentration profiles for both chromium, hollow data points, and for vanadium, filled dat~ points, respectively after heat treatment for one hour at 600°C, 700°C, 800°C, ~ ° C and lt;O0 ° C. These results were obtained by Auger spectroscopy combined with sputter profiling.
standards (dashed lines), and for the chromium-vanadium sample annealed at 600 °C for one hour (continuous lines). Indeed for this sample there are two curves in fig. 2, one corresponding to the spectrum obtained in the middle of the vanadium part of the sample, the second in the middle of the chromium part. The almost total similarity between these and the curves for the standard samples shows that although the silicides have formed, there is no mixture of the metal atoms at this stage of the annealing process. The calculated concentrations for samples annealed at 600°C (as in fig. 1), 700°C and 800°C, 900°C and 1000°C respectively, are shown in fig. 3. The data for samples heat-treated at 600 ° C and 700 ° C are barely distinguishable, but some mbdng of the metals is observed for the sample annealed at 800 ° C. Homogeneizafion of the two layers is obtained only after heat treatment at 900°C; this is particularly noticeable in the surface, vanadium-rich part of the sample. These observations based on Auger analysis are nicely corroborated from a very careful study of the corresponding backscattering spectra (not displayed here). Two sets of changes occur. (1) The bump in the spectrum for the sample annealed at 600°C, at about 1.55 MeV, that is due to the overlap of the chromium and vanadium parts of the spectrum, remains quite visible even for the sample annealed at 800 o C, but has disappeared as a result of homogeneization (and some sample roughening) after heat treatment at 900°C. (2) On the high energy end of the spectrum at about 1.68 MeV, as the annealing temperature increases, one sees the spectra moving towards slightly higher energy because chromium atoms, with their slightly higher spedfic weight, reach the surface of the sample (the initial stage of this process becomes detectible for a sample annealed at 800 ° C). The sensitivity here is limited by the stability of the accelerator and the detecting equipment; otherwise analysis
110
T.J. Finstad et al. /
is enhanced by the fact that the., spectra regardless of any sample
lyers with chromium disilicide
ace always appears fiat in backscattering ,~hness.
3, Di~usslon hbors in the fourth period of the periodic Las transition elements causes them to be cted otherwise. Thus they share the same ave nearly the same lattice [21] parame;ir melting points are quite close, 2175 o C lium and chromium. Not only are they the equilibrium diagram [22] shows that luite close, indicating that any chemical ~ts is small. The mixture is slightly ex~/ value of - 7 8 0 cal/mol. With silicon (not quite) identical series of compounds ed in table 1. The periodic table contains elements, e.g. cobalt and nickel [19], or ~'ourse the whole series of the rare-earths. ]i form a continuous [23,24] series of solid ~r the two disilieides (at 1500 ° C). [y of disilicides from TiSi 2 to WSi 2, share anit cells belong to different lattice types), have been demonstrated for some of these should be also pertain to the others. This is expected to hold for the great
Vanadium and chromium are r table of the elements. Their condi much more alike than would be e body-centered-cubic structure an ters, namely 3.0338 and 2.878/~. and 21300C, respectively for vr totally soluble into one another, the iiquidus and solidus lines al interaction between the two ele otherraic [22], with a maximum these two elements form two aim (see ref. [18]), most of which are other nearly identical neighbori~ tantalum and tungsten [15], and The metal rich silieides V.aSi and q solutions, and this is true [25] als As was pointed out [26], the fi the same basic structure (even if I
Table 1 Tile similaritiesbetween the silicidesof vanadium and chromium Composition
Structure
Latticeparameler (~) 4.722
Melting point ( ° c) 1935
V3Si
Cubic
V~Si~
Tetrahedral
9.429 4.756
2010
VSi2
Hexagonal
4.571 6.372
1670
Cr3Si
Cubic
4.564
1770
CrsSi ~
Telrahedral
9.170 4.636
1710
CrSi~
Hexagonal
4.428 6.363
1500
T.J. Finstad et a£ / Bilayers with chromium disilicide
111
dominance of the mobility of the silicon atoms over that of the metal atoms that has been proven [11-13] with bilayer experiments for MoSi2, WSi2, N'bSi 2 and TaSi 2. However, while the experiments with pairs of elements from the fourth and fifth periods present no intrinsic difficulties, there exists a certain degree of uncertainty with pairs such as niobium-vanadium or tantao lure-vanadium [13], or chromium-molybdenum,because atoms from elements of the third period are distinctly smaller than those of elements belonging to the same column but to the fourth and fifth periods; it is helpful in this respect to think of copper on the one hand, and of silver and gold on the other. While this difference in size does not, at least in the case of the last set of examples (see e.g. ref. [27]), affect greatly the diffusion coefficient, the same may not necessarily obtain for chromium and molybdenum in their respective silicides. Hence, the interest here in studying dements with nearly the same atomic sizes, chromium and vanadium (an added advantage is the near equality of the atomic weights m so that the frequency fa .~ors, proportional to l ~ m , in the diffusion coefficients should be almost equal) Thus because of the similarities between the metal atoms and their silicides one would expect that the diffusion behaviors of the two elements should be nearly the same, as if one were dealing with two isotopes. As shown in fig. 1, the disilicides form without any evidence of mixing between the metal atoms, thereby indicating that the silicide formation is dominated by the diffusion of the silicon atoms. Otherwise common sense dictates that, had the silicides formed by metal atom motion, there would be intermixing at the interface between the two layers since the two different metals should have approximately the same diffusion coefficients. Intermixing on a depth scale comparable to the thickness of the silicides occurs only after a one hour anneal at 900 ° C. Thus one may estimate that at this temperature the chromium atoms in VSi2 an~ the vanadium atoms in CrSi= have about the same mobility as the silicon atoms assumed respectively at 600 o C and 450 o C. Extrapolating and neglecting preexponential factors, that should remain of the order of 1 cm2 s - I the activation energies for diffusion should scale as the absolute temperatures, namely in CrSi 2 the ratios of activation energies, (E,), E c r / E s i and E v / E s i should be clo e to 1.4. This is the same as what is found [13] for MoSi 2 and WSi 2. One notes that the interpretation of profiles showing no metal interdiffusion presents no ambiguity, while with a one step anneal one is left uncertain about the profiles showing interpenetration; such interpenetration could have occurred in the initial metal films prior to silicidation. In ~rder to verify !b_ispnint two samples were annealed first for one hour at 550~C to form the silicides, they were then given a second hea~ treatment, respectively at 800 ° C, a~,,d 900 °C for one hour. The profiles obtained with these two samples were in no way different from those derived from samptes that received only one heat treatment at the same high temperatures. Thus, in all cases silicidation indeed occurred before metal atom mixing.
112
7"..1. Finstad et al. / Bilayers with chromium disilicide
It is clear that one did not discuss here the individual diffusion coefficients for Cr and V as would be determined by radioactive tracer measurements. The common mixing diffusion coefficient D, that may be called "diffusivity", or "chemical" diffusion coefficient, is related [21] to the individual coefficients by the relation: D = c l D 2 + c 2 D 1.
(1)
All of these anticipated to vary with concentration. Here of necessity, as in layer growth experiments [13], one must deal with average values, and one would be obliged to take the two concentrations as equal to 0.5. From then on one could estimate the values of the activation energies for diffusion in CrSi2, and VSi2, as proportional to the respective melting temperatures in kelvin. The point to be retained is that, since these melting temperatures are so close (table 1), the activation energies would be similarly close, and the two diffusion coefficients may be assumed equal within a factor of 10 in the temperature range of interest here. Diffusion should be more rapid in CrSi 2 with its slightly lower melting point. This is true for Si motion, as evidenced by the relative silicide formation temperatures. For the metals this maybe seen in fig. 3 if one looks at the V and Cr proqles for a sample annealed at 800 o C for one hour. One more problem should retain our attention: the Auger profile in fig. 3 for the sample annealed at 8 0 0 ° C for one hour appears remarkably flat, implying the possibility that each one of the two layers had become homogeneous, which would indicate that at low temperatures the two silicides may not be totally soluble. This is unlikely considering that the a values for the unit cells vary by less than 4% while the c values are nearly identical. An attempt was made to explore that question through the examination of the diffraction patterns of samples annealed at 800 ° C for various amounts of time. However, this remained fruitless because of the extremely low diffusion coefficient of the metal atoms. Results obtained at 900 ° C and 1100" C, presented below, imply that there is no phase separation in the pseudobinary system VSi2-CrSi2. In fig. 4 one may .see superimposed vertically partial diffraction patterns of samples annealed at 9 0 0 ° C for 1 h corresponding from the top to CrSi.,, VSi2, and (Cr-V)Si 2. Continuing below one finds also the diffraction patterns for similar samples annealed at 9 0 0 ° C for 15 h and 1100°C for 1 h. Because of the slightly different c / a ratios, the (003) and (111) lines for CrSi2 and VSi 2 are inverted with respect to increasing Bragg angles. There is a slight change in preferred orientations between the Cr and the V samples so that in this latter case the (003) line is absent. For a (Cr-V)Si 2 sample these two lines should be nearly superimposed. The pattern for the (Cr-V) sample annealed at 9 0 0 ° C for 1 h corresponds clearly to a two-phase material, one rich in Cr, the other rich in V; but after continued heat treatment, up to 15 h at the same temperature, the sample had evolved almost completely to the condition
T.J. Finstad et al. / Bilayers with chromium disilicide
113
o b t a i n e d a f t e r a n n e a l i n g a t 1100 ° C for 1 h, t h a t is c l e a r l y r e p r e s e n t a t i v e o f a s i n g l e p h a s e silicide. T o c o n f i r m this p o i n t o n e m a y t u r n to t a b l e 2 w h e r e the d i f f r a c t i o n p a t t e r n o b t a i n e d w i t h th is s a m p l e h as b e e n m a t c h e d a g a i n s t a c a l c u l a t e d one. F o r this p u r p o s e it w a s a s s u m e d t h a t V e g a r d ' s l a w a p p l i e s to t he se silicides, a n d t h a t the f i na l c o m p o s i t i o n c o r r e s p o n d s to th e n o n f i n a l t h i c k n e s s e s of the i n i t i a l m e t a l surfaces; for e q u a l t h i c k n e s s e s , th is gives 54 at% c h r o m i u m a n d 46 at% v a n a d i u m . T h e a g r e e m e n t b e t w e e n e × p e f i m e n t a i l a t t i c e s p a c i n g s a n d t he c a l c u l a t e d o n e s is q u i t e close. O t h e r c o n c e n t r a t i o n profiles, o b t a i n e d a ft e r a n n e a l i n g a t d i f f e r e n t t e m p e r a t u r e s a n d w i t h d i f f e r e n t a n n e a l Table 2 Diffraction results for (Vo.~Cro.s4)Si2; calculated and for sample annealed at 1100°C for 1 h hkl
Calculated o
Experimental
d (~,)
~/Io b~
d (~.)
100 101 102 110 003 111 200 103 201 112 113 104 203 211 212 114 300 301 105 213 302 220
3.892 3.321 2.465 2,250 2.123 2.119 1.946 1.864 1.861 1.836 1.543 1.474 1.435 1.433 1,335 1.299 1.297 1.271 1.211 1.209 1.201 1.123
4 15 11 19 34 100 7 1 <1 65 11 2 4 3 2 20 5 21 <1 <1 7 6
3.887 3.314
I0 10
2.247
40
2.115 1.945
150 20
1.832 1.539 1.,o,,71 1.429 1.427
20 2 2 3 3
115
1.108
8
303 214 310 205 311 312 304
1.107 1.081 1.079 1.066 1.064 1.022 1.006
2 <1 <1 1 <1 <1 2
a~ Hexagonal unit cell, a = 4.494 ,~ and c ~ 6.370 ,~. b) Relative intensifies as for CrSi 2, t e l (28]. o Peak intensity, not integrated intensity.
1o
1.296
6
1.271
30
1.200 1.125
8 14
1.105
3
114
T.J. Finstad et al. / Bilayers with chromium disilicide
003 Ill I10
~ 0~112
III
CrS~z
vsiz
900-15
,,o.,
20
22 Z4 BRAGGANGLE (8}
~6
Fig. 4. Partial diffraction patterns for CrSi2, VSi2, and for C r - V samples annealed at 900 ° C for one and 15 h, and at l l 0 0 ° C for I h.
ing times, also displayed horizontal plateaus; but the concentrations tend to increase with annealing times. It is quite likely that this is simply due to the .arge difference in the metal diffusion coefficients between lattice and grain boundaries. Rapid penetration along the grain boundaries causes these to act as secondary sources for lattice diffusion into the grains; if this is sufficiently slow the gradient along the boundaries remains negligible, yielding relatively flat diffusion profiles. Similar observations [29] were made in studying the diffusion of dopants in TiSi2.
4. Conclusions (1) An experiment with a double layer of chromium and vanadium on a silicon substrate has allowed one to show that the disilicides are formed by a process dominated by the diffusion of the silicon atoms. (2) The metal atoms are much less mobile. Their diffusion is characterized by an activation energy at least 1.4 times as high as obtains for silicon. (3) With respect to the initial question, namely the mobility of the chromium atoms during the oxidation of CrSi 2, one must make a guarded answer. At 700 o C, where the formation of pure SiO2 has been observed (e.g. ref. [30]), the chromium atoms hardly seem to be mobile enough to allow oxidation to
T.J. Finstad et aL / Bilayers with chromium disilicide
115
p r o c e e d via the reverse ( f r o m t h e silicide-oxide to the silicide-silicon interface) m o t i o n of t h e c h r o m i u m a t o m s . However, o n e k n o w s t h a t s u c h o x i d a t i o n a n d silicidation are a c c o m p a n i e d b y t h e injection of p o i n t defects. T h e s e m a y p r o v i d e t h e e x t r a mobility required to m a k e this p r o c e s s possible. (4) A s a n t i c i p a t e d t h e two disilicides a p p e a r to be c o m p l e t e l y soluble at least d o w n to 9 0 0 ° C .
Acknowledgements T h e a u t h o r s are particularly i n d e b t e d to t h e technical U n i v e r s i t y o f N o r way, in T r o n d h e i m , for t h e p e r m i s s i o n to u s e their A u g e r s p e c t r o g r a p h , to T. G a l l o w h o p r e p a r e d the v a r i o u s thin films, a n d to G. C o l e m a n w h o p e r f o r m e d s o m e o f t h e b a e k s c a t t e r i n g analyses. C o n v e r s a t i o n s w i t h P. G a s were p a r t i c u larly helpful in u n d e r s t a n d i n g t h e flat d i f f u s i o n profiles.
References [1] F.M. d'Hcorle, A. Cros, R.D. Frampton and E.A. Irene, Phil. Mag. B 55 (1987) 291. [2] H. Jiang, C.S. Petersson and M.-A. Nicolet, Thin Solid Films 140 (1986) 115. [3] C.-D. Lien, M. Bartur and M.-A. Nicolet, Mater. Res. Soc. Proc. 25 (1984) 51. [4,] J. Hudner, H. Jiang and C.S. Petersson, Le Vide, Les Couches Minces, 42-236 (1987) 63. [5] F.M. d'Heurle, Thin Solid Films 105 (1983) 285. [6] B. Leroy, Phil. Mag. B 55 (1987) 159. [7] S.M. Hu, J. Appl. Phys. 57 (1985) 4527. [81 T.Y. Tan, F. Morehead and U. G0sele, Defects in Silicon (The Electrochemical Society, Pennington, NJ, 1983) p. 325. [9] F. Brown and W.D. Mackintosh, J. Electrochem. Soc. 120 (1973) 1096. [10] W.K. Chu, S.S. Lau, J.W. Mayer, H. Miiller and K.N. Tu, Thin Solid Films 25 (1975) 393. [11] J.E. Baglin, F.M. d'Heurle, W. Hammer, C.S. Petersson and C. Serrano, J. Electron. Mater. 8 (1979) 641. [12] J.E. Baglin, F.M. d'Hcorle, W. Hammer and C.S. Petersson, Nucl. Instr. Methods 168 (1980) 491. [13] F.M. d'Heurle and P. Gas, J. Mater. Res. 1 (1986) 205. [14] A. Appelbaum, M. Eizenberg and R. Brener, Vacuum 33 0983) 227. [15] O. Thomas, F.M. d'Heurle and A. Charai, Phil. Mag. B 58 (1988) 529. [16] J.O. Olowolafe, M.-A. Nicolet and J.W. Mayer, J. Appl. Phys. 47 (1976) 5182. [17] B.I. Fomin, A.E. Gershinskii, E.I. Cherepan and F.L. Edelman, Phys. Status Solidi (a) 36 (1976) K89. [18] M.-A. Nicolet and S.S. Lau, in: VLS! Electronics, Microstructure Science, Eds. N.G. Einspruch and G.B. Larrabee (Academic Press, New York, 1983) pp. 349, 350. [19] T.J. Finstad, D.D. Anfiteatro, V.R. Deline, F.M. d'Heurle, P. Gas, V.L. Moruzzi, K. Schwartz and J. Tersoff, Thin Solid Films 135 (1986) 229. [20] L.E. Davis, N.C. MacDonald, P.W. Palmberg, G.E. Riach and R.E. Weber, Eds., Handbook of Auger Electron Spectroscopy, (Perkin Elmer, Eden Prairie. MN, 1978) pp. 75-77. [21] L.S. Darken and R,W. Gun'y, Physical Chemistry of Metals (McGraw-Hill, New York, 1953) pp. 53, 457-463.
I16
72J. Finstad et al. / Bilal r~" with chromium disilicide
[22] R. Hultgren, P. Desai, D. Hawkins, M. Gleiser and K. Kelley, Selected Values of the Thermodynamic Properties of Binary Allo3,- (American Society for Metals, Metals Park, OH, 1973) pp. 720-723. [23t W.B. Pearson, A Handbook of Lattice Spacings and Structures of Metals and Alloys (Pergamon, Oxford, 1958) p. 563. [24] H. Nowomy, H. Schroth, R. Kieffer and E Benesovsky, Monatsh. Chem. 84 (1953) 579. [25] A. Wittman, K.O. Burger and H. Nowotny. Monatsh. Chem. 93 (1962) 517. 1261 F.M. d'Heurle, in: VLSI Science and Technology/ 1982, Eds. C. Dell'Oca and W.M. Bullis (The Electrochemical Society, Pennington, N J, 1982) p. 194. [27] Y. Adda and J. Philibert, La Diffusion da~.s les Solides (Presses Universitalres de France, Paris, 1966) pp. 491-492. [28] Standard Diffraction Powder Paaern No. 35-781 (1985). [29] P. Gas, G. Scilla, A. Michel, F.K. LeGoues. O. Thomas and F.M. d'Heurle, J. Appl. Phys. 63 (1988) 5335. [30] R.D. Frampton, E.A. Irene and F.M. d'Heurle, J. Appl. Phys. 62 (1987) 2972.
IS~LAYERS WITH CHROMHJM
Applied Surface Science38 (1989) 106-116 North-Holland, Amsterdam
DISHL~CIIDE: CHROM~UM-VANADIIUM
T.L FINSTAD ~), O. THOMAS b.c) and F.M. D ' H E U R L E b.d) ~ Fystk Institutt, Universitetet i Oslo, Boks 1048, Blindern, 0316 Oslo 3, Norway bj I B M T.J. Watson Research Center. P O 218, Yorktown Heights, N Y 10598, USA ~) Laboratoire des Mat~riaux et de G~nie Physique, E N S P G , BP 46, 38402 Saint Martin d'Hdres, France •o Fasta Tillstdmdets Elektronik, K T H - Electrum, Box 1298, 164 28 Kista, Sweden
Received 20 March 1989; accepted for publication 27 April 1989
The behavior of chromium-vanadiumlayers deposited in that order on siliconsubstrates has been studied, as a function of heat treatment, with the aim of determining the difference of mobilitiesbetweenthe silicon and the metal atom.~in the two respectivedisilicides.The motion of the silicon atoms requires temperatures of the order of 500°C, that of the metal atoms temperatures about 300 °C higher.The two disilicides,previouslyreported to be totally soluble at 1500°C, appear to remain totally soluble at least down to 900°C.
I. lIntreduetion The thermal oxidation (meaning here the formation of a layer of pure SiO 2) of silicides is interesting [1-5] not only in itself, but also because it may throw light upon the still unresolved complexities of the oxidation of silicon. One aspect of this problem is the injection [6-8] of defects in the silicon during oxidation. Intriguing observations made on the oxidation of silicides seem to bear on this question. Duriog the oxidation of CoSi 2 and NiSi 2 one observes not only the decomposition [3,5] of the disilicides accompanied by the reverse motion (from the silicide-oxide interface to the silicon substrate) of the metal atoms, as during ;he formation of the disilicides, but also the motion in the same direction of about 25% of the silicon atoms which could have been oxidized. The contrast [3,4] between the formation of CrSi 2, whzn the silicon atoms constitute the most mobile species, and the oxidation of the same silicide, which occurs via the reverse motion of the chromium atoms, is equally surprising. The present experiments were undertaken in order to investigate the difference between the mobilities of the metal and the silicon atoms in CrSi 2. The experiments which have been mentioned [3,4] above were conducted with inert markers [9,10] according to an old and well established technique. This provides unambiguous but not so accurate results since they prove that CrSi 2 0169-4332/89/$03.50 © Elsevier Science Publishers B.V. (North-Holland Physics Publishing Division)
T.J. Finstad et al. / Bilayers with chrom m disihcide
107
forms predominantly by silicon motion, but ~ 'e may only say that the mobility of the metal atoms is less than about 1( + that of the silicon atoms. However, the ratio could be much smaller. In vie,,, of the suprising behaviors reported above, it becomes important to evaluate :ow big the difference might be. Namely, one wants to know as quantitatively ~ possible how different are the behaviors of the moving species during silicice formation and oxidation. To this end one proposes to study the rate at whicL similar silicides, here CrSi z and VSi 2 (one may also study pairs composed of CrSi 2 and MoSi z or W S i , ) in intimate contact with one another intermix. Th~s technique has been used to investigate the formations [11,12] of MoSi 2 and WSi z, as well as NbSi z and TaSi z. A detailed discussion is found in ref. [23]. For reasons which will become apparent later on the experiments reported here with vanadium-chromium bilayers are quite similar to experiments [14] conducted with tantalum-tungsten bilayers. In a slightly different spirit one may look also at oxidation studies carried out with niobium-tantalum bilayers [15].
2. Experimental procedures and results In order to minimize the formation of any interface oxide the two layers of chromium and vanadium were deposited in tiffs order in a single pump-down, in a vacuum chamber equipped with dual e-beam heated sources. The pressure during deposition was 1 × 10 -7 Toil', the substrate temperature 150°C: the rate of deposition 10 A / s , and the nominal film thicknesses 90 nm each (corresponding to respective disilicide thicknesses of 270 and 243 rim). The samples about 1 cm square cut from the 6 cm wafers were heat-treated in an atmosphere of helium purified over a titanium bed at about 8 5 0 ° C and zirconium at 650°C. Fig. 1, which displays several backscattering spectra, shows that the formation of CrSi2 is complete after annealing at 450 ° C for one hour, however, the formation of VSi2 requires a temperature of 600 o C. This is quite a good agreement with previously reported [16-18] results. The sensitivity of backscattering is limited for distinguishing elements with close atomic numbers such as chromium and vanadium. Thus in order to follow the effects of annealing at high temperatures one has to have recourse to Auger spectroscopy combined with sputter profiling, as was done with tantalumtungsten [14], and cobalt-nickel [19]. Supplemental information was obtained, when necessary, from X-ray diffraction patterns from a computer-controlled diffractometer equipped with a copper tube and a post-sample monochromator (which eliminates fluorescence radiations). Separate Auger spectra were obtained from CrSi 2 and VSi2, which were used as calibration standards. Three lines [20] were used for estimating the concentrations: C r L M M at about 524 eV, V L M M at about 469 eV, and Si KLL at about 1608 eV. Two of these are shawn in fig. 2, both for the
108
T.J. Finstad et aL
8/--
'
/
Bilayers with chromium disilicide
I
~
' Cr
A
V
I!0 f.4 eACKSCATTERING ENERGY (MoV)
1.8
Fig. l. Backscattering spectra of the chromlum-vanadium bilayers, respectively as-deposited and after heat treatments for one hour at 450 o C and 600 o C. Because the atomic number of vanadium is smaller than that of chromium (by one) the spectral part of vanadium is less high than that of chromium (both in the initial metallic slate and in the disilicide, after annealing at 600°C). For the same reason, and because the vanadium layer is found at the free surface the overlap of the chromium and vanadium counts causes a peak in the metal part of the spectra (at about 1.6 MeV initially, and at 1.5 MeV after completion of the disilicide reaction). Since the vanadium layer, both in the as-deposited state and immediately after silicidation, is found at the free surface the position of "surface vanadium" corresponds to the high energy side of the displayed spectra slightly below 1.7 MeV. Surface chromium would be found at a somewhat higher energy, with a vanadium-chromium separation of 11 keV for 2.3 MeV alpha particles; this would be hardly visible on the energy scale used here.
'
I
. . . . ,,,.%.~Z--
I
' ' ' " .~--'-'-Cr
,
I
' ' Si 2
I
.
'
'
I
'
'
1
a
i
VANADIUM
,
,
I
400
.
.
.
.
t
,
45O
.,
,
.
50O
.
.
.
.
550
ENERGY (eV) Fig. 2. Partial Auger spectra obtained from standard VSi 2 and CrSi 2 samples and from the middle of the vanadium-rich and chromium-rich section of a bilayer sample annealed at 700 ° C for one hour. The small peak showing at low energy for one sample only is due to nitrogen contamination during the Auger analysis. The shaded areas were used to calculate alloy concentrations.
109
T.J. Finstad et aL / Bilayers with chromium disilicide
soo 7oo 8oo 9oo Iooo
o.4
c c c c c
oc~ oc~ ~Tcr oct &cr
@v ov ~v mv Av
o.,~ o.t o
'
DEPTH ($FUTTERI~ TI~E)
Fig. 3. Concentration profiles for both chromium, hollow data points, and for vanadium, filled dat~ points, respectively after heat treatment for one hour at 600°C, 700°C, 800°C, ~ ° C and lt;O0 ° C. These results were obtained by Auger spectroscopy combined with sputter profiling.
standards (dashed lines), and for the chromium-vanadium sample annealed at 600 °C for one hour (continuous lines). Indeed for this sample there are two curves in fig. 2, one corresponding to the spectrum obtained in the middle of the vanadium part of the sample, the second in the middle of the chromium part. The almost total similarity between these and the curves for the standard samples shows that although the silicides have formed, there is no mixture of the metal atoms at this stage of the annealing process. The calculated concentrations for samples annealed at 600°C (as in fig. 1), 700°C and 800°C, 900°C and 1000°C respectively, are shown in fig. 3. The data for samples heat-treated at 600 ° C and 700 ° C are barely distinguishable, but some mbdng of the metals is observed for the sample annealed at 800 ° C. Homogeneizafion of the two layers is obtained only after heat treatment at 900°C; this is particularly noticeable in the surface, vanadium-rich part of the sample. These observations based on Auger analysis are nicely corroborated from a very careful study of the corresponding backscattering spectra (not displayed here). Two sets of changes occur. (1) The bump in the spectrum for the sample annealed at 600°C, at about 1.55 MeV, that is due to the overlap of the chromium and vanadium parts of the spectrum, remains quite visible even for the sample annealed at 800 o C, but has disappeared as a result of homogeneization (and some sample roughening) after heat treatment at 900°C. (2) On the high energy end of the spectrum at about 1.68 MeV, as the annealing temperature increases, one sees the spectra moving towards slightly higher energy because chromium atoms, with their slightly higher spedfic weight, reach the surface of the sample (the initial stage of this process becomes detectible for a sample annealed at 800 ° C). The sensitivity here is limited by the stability of the accelerator and the detecting equipment; otherwise analysis
110
T.J. Finstad et al. /
is enhanced by the fact that the., spectra regardless of any sample
lyers with chromium disilicide
ace always appears fiat in backscattering ,~hness.
3, Di~usslon hbors in the fourth period of the periodic Las transition elements causes them to be cted otherwise. Thus they share the same ave nearly the same lattice [21] parame;ir melting points are quite close, 2175 o C lium and chromium. Not only are they the equilibrium diagram [22] shows that luite close, indicating that any chemical ~ts is small. The mixture is slightly ex~/ value of - 7 8 0 cal/mol. With silicon (not quite) identical series of compounds ed in table 1. The periodic table contains elements, e.g. cobalt and nickel [19], or ~'ourse the whole series of the rare-earths. ]i form a continuous [23,24] series of solid ~r the two disilieides (at 1500 ° C). [y of disilicides from TiSi 2 to WSi 2, share anit cells belong to different lattice types), have been demonstrated for some of these should be also pertain to the others. This is expected to hold for the great
Vanadium and chromium are r table of the elements. Their condi much more alike than would be e body-centered-cubic structure an ters, namely 3.0338 and 2.878/~. and 21300C, respectively for vr totally soluble into one another, the iiquidus and solidus lines al interaction between the two ele otherraic [22], with a maximum these two elements form two aim (see ref. [18]), most of which are other nearly identical neighbori~ tantalum and tungsten [15], and The metal rich silieides V.aSi and q solutions, and this is true [25] als As was pointed out [26], the fi the same basic structure (even if I
Table 1 Tile similaritiesbetween the silicidesof vanadium and chromium Composition
Structure
Latticeparameler (~) 4.722
Melting point ( ° c) 1935
V3Si
Cubic
V~Si~
Tetrahedral
9.429 4.756
2010
VSi2
Hexagonal
4.571 6.372
1670
Cr3Si
Cubic
4.564
1770
CrsSi ~
Telrahedral
9.170 4.636
1710
CrSi~
Hexagonal
4.428 6.363
1500
T.J. Finstad et a£ / Bilayers with chromium disilicide
111
dominance of the mobility of the silicon atoms over that of the metal atoms that has been proven [11-13] with bilayer experiments for MoSi2, WSi2, N'bSi 2 and TaSi 2. However, while the experiments with pairs of elements from the fourth and fifth periods present no intrinsic difficulties, there exists a certain degree of uncertainty with pairs such as niobium-vanadium or tantao lure-vanadium [13], or chromium-molybdenum,because atoms from elements of the third period are distinctly smaller than those of elements belonging to the same column but to the fourth and fifth periods; it is helpful in this respect to think of copper on the one hand, and of silver and gold on the other. While this difference in size does not, at least in the case of the last set of examples (see e.g. ref. [27]), affect greatly the diffusion coefficient, the same may not necessarily obtain for chromium and molybdenum in their respective silicides. Hence, the interest here in studying dements with nearly the same atomic sizes, chromium and vanadium (an added advantage is the near equality of the atomic weights m so that the frequency fa .~ors, proportional to l ~ m , in the diffusion coefficients should be almost equal) Thus because of the similarities between the metal atoms and their silicides one would expect that the diffusion behaviors of the two elements should be nearly the same, as if one were dealing with two isotopes. As shown in fig. 1, the disilicides form without any evidence of mixing between the metal atoms, thereby indicating that the silicide formation is dominated by the diffusion of the silicon atoms. Otherwise common sense dictates that, had the silicides formed by metal atom motion, there would be intermixing at the interface between the two layers since the two different metals should have approximately the same diffusion coefficients. Intermixing on a depth scale comparable to the thickness of the silicides occurs only after a one hour anneal at 900 ° C. Thus one may estimate that at this temperature the chromium atoms in VSi2 an~ the vanadium atoms in CrSi= have about the same mobility as the silicon atoms assumed respectively at 600 o C and 450 o C. Extrapolating and neglecting preexponential factors, that should remain of the order of 1 cm2 s - I the activation energies for diffusion should scale as the absolute temperatures, namely in CrSi 2 the ratios of activation energies, (E,), E c r / E s i and E v / E s i should be clo e to 1.4. This is the same as what is found [13] for MoSi 2 and WSi 2. One notes that the interpretation of profiles showing no metal interdiffusion presents no ambiguity, while with a one step anneal one is left uncertain about the profiles showing interpenetration; such interpenetration could have occurred in the initial metal films prior to silicidation. In ~rder to verify !b_ispnint two samples were annealed first for one hour at 550~C to form the silicides, they were then given a second hea~ treatment, respectively at 800 ° C, a~,,d 900 °C for one hour. The profiles obtained with these two samples were in no way different from those derived from samptes that received only one heat treatment at the same high temperatures. Thus, in all cases silicidation indeed occurred before metal atom mixing.
112
7"..1. Finstad et al. / Bilayers with chromium disilicide
It is clear that one did not discuss here the individual diffusion coefficients for Cr and V as would be determined by radioactive tracer measurements. The common mixing diffusion coefficient D, that may be called "diffusivity", or "chemical" diffusion coefficient, is related [21] to the individual coefficients by the relation: D = c l D 2 + c 2 D 1.
(1)
All of these anticipated to vary with concentration. Here of necessity, as in layer growth experiments [13], one must deal with average values, and one would be obliged to take the two concentrations as equal to 0.5. From then on one could estimate the values of the activation energies for diffusion in CrSi2, and VSi2, as proportional to the respective melting temperatures in kelvin. The point to be retained is that, since these melting temperatures are so close (table 1), the activation energies would be similarly close, and the two diffusion coefficients may be assumed equal within a factor of 10 in the temperature range of interest here. Diffusion should be more rapid in CrSi 2 with its slightly lower melting point. This is true for Si motion, as evidenced by the relative silicide formation temperatures. For the metals this maybe seen in fig. 3 if one looks at the V and Cr proqles for a sample annealed at 800 o C for one hour. One more problem should retain our attention: the Auger profile in fig. 3 for the sample annealed at 8 0 0 ° C for one hour appears remarkably flat, implying the possibility that each one of the two layers had become homogeneous, which would indicate that at low temperatures the two silicides may not be totally soluble. This is unlikely considering that the a values for the unit cells vary by less than 4% while the c values are nearly identical. An attempt was made to explore that question through the examination of the diffraction patterns of samples annealed at 800 ° C for various amounts of time. However, this remained fruitless because of the extremely low diffusion coefficient of the metal atoms. Results obtained at 900 ° C and 1100" C, presented below, imply that there is no phase separation in the pseudobinary system VSi2-CrSi2. In fig. 4 one may .see superimposed vertically partial diffraction patterns of samples annealed at 9 0 0 ° C for 1 h corresponding from the top to CrSi.,, VSi2, and (Cr-V)Si 2. Continuing below one finds also the diffraction patterns for similar samples annealed at 9 0 0 ° C for 15 h and 1100°C for 1 h. Because of the slightly different c / a ratios, the (003) and (111) lines for CrSi2 and VSi 2 are inverted with respect to increasing Bragg angles. There is a slight change in preferred orientations between the Cr and the V samples so that in this latter case the (003) line is absent. For a (Cr-V)Si 2 sample these two lines should be nearly superimposed. The pattern for the (Cr-V) sample annealed at 9 0 0 ° C for 1 h corresponds clearly to a two-phase material, one rich in Cr, the other rich in V; but after continued heat treatment, up to 15 h at the same temperature, the sample had evolved almost completely to the condition
T.J. Finstad et al. / Bilayers with chromium disilicide
113
o b t a i n e d a f t e r a n n e a l i n g a t 1100 ° C for 1 h, t h a t is c l e a r l y r e p r e s e n t a t i v e o f a s i n g l e p h a s e silicide. T o c o n f i r m this p o i n t o n e m a y t u r n to t a b l e 2 w h e r e the d i f f r a c t i o n p a t t e r n o b t a i n e d w i t h th is s a m p l e h as b e e n m a t c h e d a g a i n s t a c a l c u l a t e d one. F o r this p u r p o s e it w a s a s s u m e d t h a t V e g a r d ' s l a w a p p l i e s to t he se silicides, a n d t h a t the f i na l c o m p o s i t i o n c o r r e s p o n d s to th e n o n f i n a l t h i c k n e s s e s of the i n i t i a l m e t a l surfaces; for e q u a l t h i c k n e s s e s , th is gives 54 at% c h r o m i u m a n d 46 at% v a n a d i u m . T h e a g r e e m e n t b e t w e e n e × p e f i m e n t a i l a t t i c e s p a c i n g s a n d t he c a l c u l a t e d o n e s is q u i t e close. O t h e r c o n c e n t r a t i o n profiles, o b t a i n e d a ft e r a n n e a l i n g a t d i f f e r e n t t e m p e r a t u r e s a n d w i t h d i f f e r e n t a n n e a l Table 2 Diffraction results for (Vo.~Cro.s4)Si2; calculated and for sample annealed at 1100°C for 1 h hkl
Calculated o
Experimental
d (~,)
~/Io b~
d (~.)
100 101 102 110 003 111 200 103 201 112 113 104 203 211 212 114 300 301 105 213 302 220
3.892 3.321 2.465 2,250 2.123 2.119 1.946 1.864 1.861 1.836 1.543 1.474 1.435 1.433 1,335 1.299 1.297 1.271 1.211 1.209 1.201 1.123
4 15 11 19 34 100 7 1 <1 65 11 2 4 3 2 20 5 21 <1 <1 7 6
3.887 3.314
I0 10
2.247
40
2.115 1.945
150 20
1.832 1.539 1.,o,,71 1.429 1.427
20 2 2 3 3
115
1.108
8
303 214 310 205 311 312 304
1.107 1.081 1.079 1.066 1.064 1.022 1.006
2 <1 <1 1 <1 <1 2
a~ Hexagonal unit cell, a = 4.494 ,~ and c ~ 6.370 ,~. b) Relative intensifies as for CrSi 2, t e l (28]. o Peak intensity, not integrated intensity.
1o
1.296
6
1.271
30
1.200 1.125
8 14
1.105
3
114
T.J. Finstad et al. / Bilayers with chromium disilicide
003 Ill I10
~ 0~112
III
CrS~z
vsiz
900-15
,,o.,
20
22 Z4 BRAGGANGLE (8}
~6
Fig. 4. Partial diffraction patterns for CrSi2, VSi2, and for C r - V samples annealed at 900 ° C for one and 15 h, and at l l 0 0 ° C for I h.
ing times, also displayed horizontal plateaus; but the concentrations tend to increase with annealing times. It is quite likely that this is simply due to the .arge difference in the metal diffusion coefficients between lattice and grain boundaries. Rapid penetration along the grain boundaries causes these to act as secondary sources for lattice diffusion into the grains; if this is sufficiently slow the gradient along the boundaries remains negligible, yielding relatively flat diffusion profiles. Similar observations [29] were made in studying the diffusion of dopants in TiSi2.
4. Conclusions (1) An experiment with a double layer of chromium and vanadium on a silicon substrate has allowed one to show that the disilicides are formed by a process dominated by the diffusion of the silicon atoms. (2) The metal atoms are much less mobile. Their diffusion is characterized by an activation energy at least 1.4 times as high as obtains for silicon. (3) With respect to the initial question, namely the mobility of the chromium atoms during the oxidation of CrSi 2, one must make a guarded answer. At 700 o C, where the formation of pure SiO2 has been observed (e.g. ref. [30]), the chromium atoms hardly seem to be mobile enough to allow oxidation to
T.J. Finstad et aL / Bilayers with chromium disilicide
115
p r o c e e d via the reverse ( f r o m t h e silicide-oxide to the silicide-silicon interface) m o t i o n of t h e c h r o m i u m a t o m s . However, o n e k n o w s t h a t s u c h o x i d a t i o n a n d silicidation are a c c o m p a n i e d b y t h e injection of p o i n t defects. T h e s e m a y p r o v i d e t h e e x t r a mobility required to m a k e this p r o c e s s possible. (4) A s a n t i c i p a t e d t h e two disilicides a p p e a r to be c o m p l e t e l y soluble at least d o w n to 9 0 0 ° C .
Acknowledgements T h e a u t h o r s are particularly i n d e b t e d to t h e technical U n i v e r s i t y o f N o r way, in T r o n d h e i m , for t h e p e r m i s s i o n to u s e their A u g e r s p e c t r o g r a p h , to T. G a l l o w h o p r e p a r e d the v a r i o u s thin films, a n d to G. C o l e m a n w h o p e r f o r m e d s o m e o f t h e b a e k s c a t t e r i n g analyses. C o n v e r s a t i o n s w i t h P. G a s were p a r t i c u larly helpful in u n d e r s t a n d i n g t h e flat d i f f u s i o n profiles.
References [1] F.M. d'Hcorle, A. Cros, R.D. Frampton and E.A. Irene, Phil. Mag. B 55 (1987) 291. [2] H. Jiang, C.S. Petersson and M.-A. Nicolet, Thin Solid Films 140 (1986) 115. [3] C.-D. Lien, M. Bartur and M.-A. Nicolet, Mater. Res. Soc. Proc. 25 (1984) 51. [4,] J. Hudner, H. Jiang and C.S. Petersson, Le Vide, Les Couches Minces, 42-236 (1987) 63. [5] F.M. d'Heurle, Thin Solid Films 105 (1983) 285. [6] B. Leroy, Phil. Mag. B 55 (1987) 159. [7] S.M. Hu, J. Appl. Phys. 57 (1985) 4527. [81 T.Y. Tan, F. Morehead and U. G0sele, Defects in Silicon (The Electrochemical Society, Pennington, NJ, 1983) p. 325. [9] F. Brown and W.D. Mackintosh, J. Electrochem. Soc. 120 (1973) 1096. [10] W.K. Chu, S.S. Lau, J.W. Mayer, H. Miiller and K.N. Tu, Thin Solid Films 25 (1975) 393. [11] J.E. Baglin, F.M. d'Heurle, W. Hammer, C.S. Petersson and C. Serrano, J. Electron. Mater. 8 (1979) 641. [12] J.E. Baglin, F.M. d'Hcorle, W. Hammer and C.S. Petersson, Nucl. Instr. Methods 168 (1980) 491. [13] F.M. d'Heurle and P. Gas, J. Mater. Res. 1 (1986) 205. [14] A. Appelbaum, M. Eizenberg and R. Brener, Vacuum 33 0983) 227. [15] O. Thomas, F.M. d'Heurle and A. Charai, Phil. Mag. B 58 (1988) 529. [16] J.O. Olowolafe, M.-A. Nicolet and J.W. Mayer, J. Appl. Phys. 47 (1976) 5182. [17] B.I. Fomin, A.E. Gershinskii, E.I. Cherepan and F.L. Edelman, Phys. Status Solidi (a) 36 (1976) K89. [18] M.-A. Nicolet and S.S. Lau, in: VLS! Electronics, Microstructure Science, Eds. N.G. Einspruch and G.B. Larrabee (Academic Press, New York, 1983) pp. 349, 350. [19] T.J. Finstad, D.D. Anfiteatro, V.R. Deline, F.M. d'Heurle, P. Gas, V.L. Moruzzi, K. Schwartz and J. Tersoff, Thin Solid Films 135 (1986) 229. [20] L.E. Davis, N.C. MacDonald, P.W. Palmberg, G.E. Riach and R.E. Weber, Eds., Handbook of Auger Electron Spectroscopy, (Perkin Elmer, Eden Prairie. MN, 1978) pp. 75-77. [21] L.S. Darken and R,W. Gun'y, Physical Chemistry of Metals (McGraw-Hill, New York, 1953) pp. 53, 457-463.
I16
72J. Finstad et al. / Bilal r~" with chromium disilicide
[22] R. Hultgren, P. Desai, D. Hawkins, M. Gleiser and K. Kelley, Selected Values of the Thermodynamic Properties of Binary Allo3,- (American Society for Metals, Metals Park, OH, 1973) pp. 720-723. [23t W.B. Pearson, A Handbook of Lattice Spacings and Structures of Metals and Alloys (Pergamon, Oxford, 1958) p. 563. [24] H. Nowomy, H. Schroth, R. Kieffer and E Benesovsky, Monatsh. Chem. 84 (1953) 579. [25] A. Wittman, K.O. Burger and H. Nowotny. Monatsh. Chem. 93 (1962) 517. 1261 F.M. d'Heurle, in: VLSI Science and Technology/ 1982, Eds. C. Dell'Oca and W.M. Bullis (The Electrochemical Society, Pennington, N J, 1982) p. 194. [27] Y. Adda and J. Philibert, La Diffusion da~.s les Solides (Presses Universitalres de France, Paris, 1966) pp. 491-492. [28] Standard Diffraction Powder Paaern No. 35-781 (1985). [29] P. Gas, G. Scilla, A. Michel, F.K. LeGoues. O. Thomas and F.M. d'Heurle, J. Appl. Phys. 63 (1988) 5335. [30] R.D. Frampton, E.A. Irene and F.M. d'Heurle, J. Appl. Phys. 62 (1987) 2972.