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Rapid formation of intermetallic compounds interdiffusion in the CuSn and NiSn systems
Acta metall, mater. Vol. 43, No. 1, pp. 329-337, 1995
~
Pergamon
Copyright © 1994ElsevierScienceLtd Printed in Great Britain.All rights reserved 0956-7151/94$7.00+ 0.00
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RAPID FORMATION OF INTERMETALLIC C O M P O U N D S BY I N T E R D I F F U S I O N IN THE Cu-Sn A N D Ni-Sn SYSTEMS S. BADER1, W. GUST1 and H. HIEBER2 JMax-Planck-Institut ffir Metallforschung and Institut f/ir Metallkunde, Seestr. 75, D-70174 Stuttgart and 2Centrum fiir Mikroverbindungstechnik in der Elektronik, Forschung und Entwicklung GmbH, Ilsahl 5, D-24503 Neumiinster, Germany (Received 9 March 1994)
Abstract--Fundamental investigations were carried out to examine a new interconnection technology which is based on the rapid formation of intermetallic compounds composed of a high melting component (e.g. Cu or Ni) and a low melting component (e.g. Sn) between two layers of the high melting component at temperatures just above the melting point of Sn. This reaction is known as isothermal solidification. The growth of the intermetallic compounds of the Cu-Sn and Ni-Sn systems were studied in thin films at temperatures from 513 to 673 K. At the beginning of interdiffusion, the Sn-rich intermetallic compounds r/ (Cu6Sn5) and Ni3Sn4 grow fastest with a non-parabolic time dependence. The Cu3Sn phase grows parabolically, however, not the Ni3Sn2 phase. The Ni3Sn phase does not nucleate below 623 K in specimens with clean Ni/Sn interfaces. For the first time, the influence of the small grain size of the Cu or Ni thin filmswere studied by performing similar experimentswith Cu or Ni singlecrystals. In the Cu-Sn system the interdiffusion coefficients for Cu3Sn obtained from the thin film experiments are twice those obtained from the single crystal experiments. In the Ni-Sn system there are no differences between thin film and single crystal results.
1. INTRODUCTION The formation and growth of intermetallics (IMs) in thin C u - S n - C u and N i - S n - N i films have been studied. The reason for these investigations was the examination of a new thin film contact technology [1-3]. This technology is based on the rapid formation of IMs composed of a high melting component (HC), such as Cu and Ni, and a low melting component (LC), such as Sn, at temperatures just above the melting point of the LC between two layers of the HC. Figure 1 shows this schematically. The bond-layer grows by liquid-solid interdiffusion at the thin film interfaces and may consist of one or more IMs. After consumption of the liquid LC the interconnection is stable up to the melting temperature, Tin, of the IM with the lowest melting point existing in the bond-layer. Figure 2 shows this for the Ni-Sn system. By variation of the film thicknesses of the HC and LC and the time of the heat treatment, the kind and proportionate amount of IMs can be influenced. In this way small and heat-resistant interconnections can be produced. In contrast to the soft-solder technology the maximal use temperature is here appreciably higher than the bonding temperature, Tb, at which the bond is produced Tm> Tb + 500 K.
(1)
This technology can therefore be applied to the packaging of devices which exhibit high heat dissi-
pation or which are exposed to high temperatures or changes in ambient temperature. The feasibility of this technology depends on the growth velocity, the aging behavior and the morphology of the IMs in the bonding layer. Initial experiments showed that all possible systems HC-LC exhibited high growth velocities for the IMs. The Cu-Sn and Ni-Sn systems were choosen for further studies because they have only a few IMs and the compositions and structures of the IMs are all well-known.
2. EXPERIMENTALPROCEDURE The specimens were made by sequential sputtering of Cu or Ni and Sn layers on Si wafers (0.35 mm thick). The specimens had dimensions of 4 x 4 × 0.35 mm 3 and were made by cleaving from a appropriately scribed Si wafer. Because of the low adhesion of Cu on Si a Ni-Cr-AI adhesion layer was used. To avoid any influence of the adhesion layer on the growth of the IMs, the Cu layers were made thicker. To produce the films, a 2400 type Perkin Elmer sputter depositer was used. The base pressure before sputtering was 5.3 x 10-4 Pa and the Ar pressure while sputtering was 1.3 Pa. The following combinations of film thicknesses were produced: Cu(7/~m)/Sn(2.6/~m), Cu(7 pm)/Sn(l.6/~m), Ni(4.4/~m)/Sn(4.3/tm) and Ni(4.1 #m)/Sn(2.6/~m). During sputtering of the Sn layers onto Cu, an
329
330
BADER et aL: FORMATION OF INTERMETALLIC COMPOUNDS
approx. 0.3 # m thick layer of IMs grew. In the Ni-Sn system there was no growth of IMs while sputtering. Sputtered layers of the HC have a very small grain size. To study the influence of this small grain size, similar experiments using Cu and Ni single crystals were done. For these experiments Sn layers comparable to the Sn layers of the sputtered specimens were evaporated onto polished single crystal specimens having the dimensions 4 x 4 x 1 mm 3. Two pieces were put together as shown in Fig. 1 and placed in a specimen holder, which exerted a small but uniform pressure on the specimen through a central screw. This "two piece" arrangement caused the heating up time to be longer because of the heat capacity of the specimen holder, however, it prevented the formation of Sn droplets while melting, enlarged the useable interface length for measuring the width of the ICs and was very similar to a later application. Heat treatments up to 20 min were carried out in a big (186 cm 3) Pb-Sn solder bath. In this case, the specimens were wrapped in A1 foil to prevent reactions with the Pb-Sn solder. The effective heating up time was only 2-3 s and could be neglected in the analysis. Longer annealings were carried out in an oil bath or in tube furnaces. Here the specimens with specimen holders were sealed in evacuated glass tubes backfilled with 200 mbar Ar to protect them against oil or corrosion. Most of the observations were made using an optical microscope (LM). Only in special cases was the scanning electron microscope (SEM) used. The annealed specimens were mounted in epoxy and ground perpendicular to the metal and IM layers with 600, 1200 and 4000 grit abrasive paper. Polishing followed with 3 and 1 # m diamond paste and colloidal SiO 2 suspension [4]. If Sn was still visible after
Before reaction
HC
HC
Sub
Sub
After reaction
SubI
t F--
5
0
0
°°°I/ 3
0 N1
~
T2
\
......:::1 0
20
0
40 at.%
~
60 Sn
80
100 Sn
Fig. 2. Ni-Sn phase diagram [9]. If there is Ni3Sn4 phase in the bond area, the bond is stable up to T~. If there is only Ni3Sn2phase in the bond area, the bond is stable up to T2. L = liquid phase, Tb = bonding temperature. the heat treatment, the IM/Sn interface could be made more visible by selectively etching away the Sn with an aqueous solution of 35 g/l ortho-nitrophenol and 50 g/1 NaOH. The ability to distinguish between the other phases could not be improved by chemical and electrochemical methods. The interference layer method [5] was the solution to this problem. Here a thin oxide film is reactively sputtered on the surface of the specimen to be used as a reflection interference filter. In this way, the bright~lark and color contrasts are improved. For the Cu-Sn specimens a Pt anode was used, and for the Ni-Sn specimens a Pb anode. For the optical microscope observations, a Zeiss Axiophot-type microscope with an oil differential interference contrast objective and a magnification of 1000:1 was used. The width of the IM layers were measured on the screen of a video system connected to the microscope at a total magnification of 5000: 1. Because some of the IM interfaces grew unevenly and did not remain planar the average thickness of an IM layer was determined by measuring the layer thickness at 50 equally spaced points and calculating the average thickness. The measurements are known to within +0.2 #m. The standard deviation for different IM layers was between 0.05 and 0.15 #m. The identification of the IMs were carried out by X-ray diffraction measurements. A Guinier camera with asymmetric transmission arrangement was used. The multilayer thin films were stripped off from their substrates and examined directly. For identification, the measured scans were compared with calculated scans for the different ICs. The data for the calculations were obtained from the literature [6]. In all cases, the IMs could be identified as the stable IMs of the phase diagrams.
HC
IM
Sub
3. RESULTS AND DISCUSSION
3.1. Growth of Sn-rich intermetallic compounds in the presence of liquid Sn C
Fig. 1. Structure of the bond area before and after heat treatment. HC, LC = high and low melting component, Sub = substrate, IM = intermetallic compound.
The IMs grow by liquid-solid interdiffusion. First, the changes in the border phases will be considered. The saturation of the Sn(L) is very quick, because the saturation concentrations of the HC in Sn(L) are low
BADER et al.:
FORMATION OF INTERMETALLIC COMPOUNDS
in the temperature range studied (513-603 K) and the diffusion in the liquid is very fast. An estimation and measurements by Tosima e t al. [7] showed that the saturation time is between 1 and 3 s. The saturation concentration of Sn in the HC is very low as well, and the measurements showed that the decrease in thickness of the HC at the HC/IM interface for the IM formation is higher than the penetration depth of Sn in the HC. So the border phases are the pure HC and saturated Sn(L). The beginning of the Cu-Sn interdiffusion was studied at 513 and 573 K using the specimens with a 2 × 2.6/~m thick Sn layer. The r/ and E phases form simultaneously [Fig. 3(a)]. The growth velocity of the r/phase is very fast compared to that of the E phase. While the E phase grows in the form of a continuous and wavy layer, the r/phase grows in a hemispherical shape [Fig. 3(a, b)]. The hemispheres were identified
as single crystals using polarized light. Between them are often deep grooves down to the Cu. At these places the thickness of the Cu- and E-layer are minimums. Therefore, there are grounds for the supposition that these grooves are important diffusion paths for Cu and Sn in addition to bulk diffusion through the ~/ phase. The Cu enters the Sn(L) at the base of the grooves and moves quickly because of the high diffusivity and convection to the top of the hemispheres. With time, the hemispheres grow perpendicular to the Cu surface and by a process comparable to grain growth parallel to it, while the hemispherical shape is preserved. The growth exponent of the grain growth is similar to that of particle coarsening and is predicted by the Lifshitz-Wagner law. The grain size at higher temperature is always larger than that at lower temperature. After 8-10min at 513 K, the Sn is changed completely into the IMs [Fig. 3(c)]. The ~/ crystals extend across the interconnection area. The grain boundaries lie perpendicular to the layers. A faultless and voidless interconnection zone has grown between the Cu layers. The ~/ crystallites extend through the entire zone from one Cu layer to the other. Defects in the interconnection zone can form in the shape of pores. They form for instance if the pressure on the specimen is too low during annealing, because of the volume contraction, which occurs during formation of the IMs and which is 7% for the Cu-Sn system. The pores often form in the shape of long thin gaps (fissures) between the r/ crystallites, when there is insufficient Sn for the crystals to grow together. The shape of the hemispheres made it difficult to measure the thickness, Ax, of the r/ phase and increased the error in these measurements. The average thickness was determined as described above. The annealing times, t, varied from 5 s to 5 min. Until the r/ phase reached a thickness of 2.0-2.5 #m, the distance between the r/ layers on both sides were large enough so they did not influence each other. The diffusion-controlled growth of an IM normally obeys a parabolic growth law [8] A x = ko " t ~
Fig. 3. 513 K, Cu/5.2/~m Sn/Cu (a-c) and Cu/3.2/~m Sn/Cu (d) after I min (a), 30s (b), 12min (c) and 20min (d). Sn selectively etched away (a, b).
331
or
l o g A x = l o g ko ÷ n l o g t
(2)
where n equals 0.5. Table 1 shows the measured widths of the ~/and E layers. The data from the 513 K experiments are the average of 2-4 measurements. In Fig. 4, the data for the ~/ phase are plotted in a log Ax-log t diagram. The n and k0 values computed from the curves in this figure by linear regression are the following: n = 0.21 and 0.25 and k 0 = 1.906 and 1.800/tm/min n for 513 and 573 K, respectively. The following two observations were noted: (a) the growth of the t/ phase does not obey a parabolic growth law at the beginning of interdiffusion; (b) at short annealing times, the average widths of the r/ layers for the 513 K experiments are larger than those for the 573 K experiments. The deviation from the parabolic growth law
332
BADER et al.:
FORMATION OF INTERMETALLIC COMPOUNDS
Table 1. Average thickness of the r/- and E-phase layers in the experiments where liquid Sn was still present Layer width (,am) Experiment t at 513 K (min) Ax, Ax~ Ax'~ 1-3 0.083 0.15 1.15 1.25 4-6 0.25 0.20 1.40 1.59 7-10 0.50 0.26 1.68 1.85 11-13 1.00 0.30 1.85 2.06 14-15 2.50 0.40 2.35 2.63
t
(min) 0.17 0.18 0.53 0.50 1.00 1.00 2.00 2.00
Ax, 0.25 0.25 0.58 0.65 0.75 0.75 0.80 0.83
Ax~ 1.27 1.18 1.62 1.47 1.82 1.74 2.12 2.17
Ax~ = Ax~ + Axe,.
(3)
In this equation Axe, is part of the q-layer thickness, which was used to form the e phase. The following reaction is responsible for the formation of the e phase 9 Cu + C u 6 Sn5 ~ 5 Cu3 Sn.
Layer width (#m) Experiment at 573 K 1 2 3 4 5 6 7 8
F o r this reason an q-layer thickness Ax~ was calculated using the equation
Ax~ 1.44 1.35 2.02 1.92 2.34 2.26 2.67 2.74
Thus the following expression for Ax,, results Ax,, = 11 V, Ax,/20 V,
(4)
where V, and V, are the molar volumes of the q and c phases, respectively. If the reduction of the q phase is taken into account by using Ax~ instead of Ax~ in Fig. 4, the point of intersection of the two lines lies
cannot be attributed only to the fact that the IM layer formed during deposition was ignored. A n estimation showed this not to be the case. The reason for the deviation is probably that a large part of transport of the H C and Sn to the phase boundaries, where the formation of the q phase takes place, runs through the grooves between q crystallites. In the course of time, these grooves get narrower and their density is reduced by a process comparable to grain growth or Ostwald ripening. This means that the diffusion slows down not only by the increase in diffusion path length and reduction in concentration gradient, but also by the reduction of transport capacity of the grooves. This is manifested in n values less than 0.5. Our explanation for the second observation is that in addition to the q phase, the e phase is also growing. As mentioned above, the q phase also grows by diffusion of copper through the liquid Sn. The activation energy for the Cu diffusion in Sn(L) is only 17.5 kJ/mol [10]. The q phase is both produced by reaction of E phase with Sn and is reduced to E phase by reaction of q phase with Cu. The activation energy for interdiffusion in the E phase is 65 kJ/mol. This means that the ratio of q growth rate is smaller at higher temperatures.
0.%
0.4 -E ._~0.3 t~0.2 0
0.1 0.0
- i .2 ]og t (min)
Fig. 4. Growth of the r/(Cu6Sns) phase at 513 and 573 K.
Fig. 5. Ni/8.6 #m Sn/Ni at 513 K after 165 s (a), 60 s (b) and 90 min (c). Sn selectively etched away (a, b); section parallel to intermetallic layers (b).
BADER et al.:
-0.1
-1.0
FORMATION OF INTERMETALLIC COMPOUNDS
5"13 K
-0.5
0.0
0.5
.0
log t (rain) Fig. " .~erease of the Ni film thickness during the reaction to form Ni3Sn4 as a function of time at 513 K.
at a time of 20 s. In addition, the saturation concentration of Cu in Sn(L) is higher at higher temperatures; therefore, more time is necessary to saturate the Sn. The lateral size of the r/ hemispheres is larger at higher temperatures, thus decreasing the total amount of fast diffusion path (grooves). Therefore, the growth rate of the r/ phase is smaller at higher temperatures in the initial stages of annealing. The initial stages of growth in the Ni-Sn system were studied by using specimens with a 2 x 4.3 p m thick Sn layer. Figure 5(a, b) shows SEM photographs perpendicular and parallel, respectively, to the layers of specimens with short annealing times. At the start of annealing, only one of the three IMs of the Ni-Sn system (Fig. 2) appears. This initial IM was identified as the Ni3Sn4 phase and was found to exhibit three different morphologies simultaneously: as a fine-grained, planar layer at the Ni interface; long, thin, idiomorphic whiskers; and large, polygonal, idiomorphic crystals. After an annealing time of only 7 s, a high density of Ni3Sn 4 whiskers is present. The large crystals are not yet seen. The density and thickness of the whiskers reach their maximum after 2 min, after which they decrease. After 30 s, the first large crystals appear and their number increases rapidly from the first to fifth minute of annealing. Obviously a recrystallization of the whiskers into the large crystals takes place. The driving force for this recrystallization is the reduction of surface area and energy of the Ni3 Sn4 phase. Meanwhile the fine-grained layer grows obeying parabolic growth law. After 60 min, the Sn is completely changed into Ni3Sn 4. The entire layer now consists of two equally thick ( ~ 1 #m) finegrained layers at the Ni interfaces and a coarse grained area in the middle [Fig. 5(c)]. When the pressure on the specimen was too low during annealing, there is a band of small pores between these two areas. The quick growth and the idiomorphic shape of the whiskers and big crystals indicate that the N i 3 S n 4 phase initially forms almost exclusively by the diffusion of Ni through the Sn(L) and crystallization from the liquid phase. The low solubility of Ni in Sn(L) and the high density of whiskers preclude AM 43/I--V
333
whiskers from forming during cooling. Obviously the NiaSn4 phase has a special lattice direction with a high crystallization velocity. Grains with the appropriate orientation grow out of the fine-grained layer and form the whiskers. The shape of the Ni 3Sn 4 phase made it impossible to measure the average layer width directly. The relative amount of Ni3Sn 4 phase formed was determined by measuring the decrease of the Ni layer thickness, ANi. Because the thickness of the Ni layer was not uniform across the Si wafer, the thickness of each specimen was measured before and after the heat treatment. Therefore the error of ANi was relatively high. Figure 6 shows the log Ax-log t plot of the measured data. When the data were fit with a straight line by linear regression the following k and n values were obtained: n = 0.28, k = 0.453 #m/min °'28. As was the case for the Cu-Sn, the growth exponent is smaller than 0.5, and the reason for this behavior is the same as that for Cu-Sn. In the course of time the transport capacity of the diffusion paths through the Sn(L) decreases because of the increasing density of whiskers and large crystals. The morphology in both systems is shaped by the fact that part of the Cu and Sn diffusion goes through the Sn(L). Because of the higher interface area, these morphologies are energetically less favorable; however, they result in an increase in the temporal reduction of free energy. Probably the morphology that is formed is mainly determined by the relation of the effective interdiffusion coefficients [see equation (8)] or transport capacities of the IMs and the Sn(L). If the effective interdiffusion coefficients, Def, of the IMs are low, then the transport through the Sn(L) predominates. This results in the rugged or open morphology and the idiomorphic shapes of the IC, which form by crystallization from the liquid phase. The melting point of Ni3 Sn4 is 1087 K, which means that, at 513 K, Oef of the Ni3Sn 4 phase is probably lower than D~f of the r/ phase which has a melting point of 688 K. Therefore the transport of the HC through the Sn(L) is probably larger in the case of Ni-Sn than in the case of Cu-Sn. This is visible in the
0 . 4 F
0.5
.......
T ........... ~
0.0
0.5 ]og
--7
1.0
]5
0
t (rnin)
Fig. 7. Growth of the E(Cu 3Sn) phase in the temperature range from 513 to 603 K.
BADER et al.: FORMATION OF INTERMETALLIC COMPOUNDS
334
600
500 K
-13
© E
Fig. 9. Cu/3.2 #m Sn/Cu, 30 min at 543 K. -15 15
1'8 i i04/T (I/K) Fig. 8. Arrhenius plot of Def of the thin film and single crystal experiments. morphology. Other factors which influence the morphology are the specific interface energy of the IM/Sn(L) boundary and the anisotropy of the effective interdiffusion and of the crystallization.
3.2. Growth o f Cu- or Ni-rich intermetallic compounds In the Cu-Sn system the growth of the e(Cu3Sn) phase was studied in the temperature range from 513 to 603 K using the specimens with 2 x 1.6 # m thick Sn layers. The Sn changes to the r/phase within a few seconds, and the slower growing E phase grows as it does in a Cu/r/(Cu6 Sn5) diffusion couple. The E phase grows in a planar shape [Fig. 3(d)]. At short annealing times and low temperatures no grain structure is visible. For longer heat treatments and higher temperatures a very fine grain structure of thin, columnar grains, which reach through the entire layer, is visible under polarized light. The lateral grain size (parallel to the layers) is less than 1 p m and was too small to be measured. The measured data for the layer thickness are plotted in Fig. 7. The E phase grows to a layer thickness of 2 # m within 40 rain at 513 K and within 4 min at 603 K. The data were fitted with equation (2), and the n values were found between 0.45 and 0.49. Therefore the parabolic growth law was assumed to be valid. The parabolic growth constants, kp
kp = Ax2/2t
To compute the effective interdiffusion coefficients, equations derived by Wagner [8] were used. For interdiffusion in the E phase in a Cu/r/(Cu6Sns) diffusion couple the following equation is valid Def = ( c , - Ccu)(C~ - _c,)kp
(6)
Ac,(c, - Cc.)
where c is the mole fraction of Sn, c, is the average composition of the E phase, Cc. is the saturation concentration of Sn in Cu (Cc. = 0), c, is the Sn concentration of the r/phase in equilibrium with the E phase and Ac, is the homogeneity range of the E phase. The concentrations are taken from the phase diagram [9]. The Def values are plotted in Fig. 8 and fitted with an Arrhenius equation
Oet= O°f exp(-- Q,f/RT)
(7)
where R is the gas constant and T is the absolute temperature. The determined values of the preexponential factor, De°, and the activation energy, Qef, are listed in Table 2. Figure 9 shows a specimen annealed for 30 min at 543 K. Both e layers have grown together. The interconnection zone is widely free of pores. To study the influence of the small grain size of the thin films, similar experiments with Cu single crystals were made. The morphology of the E layer was the same. The growth of the E phase obeys the parabolic growth law too (0.48 < n < 0.49). The calculated D~f values are plotted in Fig. 8. The following differences
(5)
determined are listed in Table 2.
Table 2. Parabolicgrowth constant (kp) and Arrheniusparameters (Deft Qef) for the Ephase kp (10 lSm2/s) T (K) Thin film Singlecrystal 513 0.92 0.38 543 2.00 0.94 573 4.64 2.85 603 9.30 -Specimens Thin film Singlecrystal
D~r(10 8m2/s) 5.3 80.8
Qef(kJ/mol) 66.1 81.6
lIT Fig. 10. Schematic description of the temperature dependence of De of the E(Cu3Sn) phase. A = temperature range studied; D0f Xr, Dsc = D,f of the thin film and single crystal experiments; D, Db = bulk and grain boundary diffusion coefficients.
BADER et al.: FORMATION OF INTERMETALLIC COMPOUNDS
335
between the thin film and single crystal experiments were observed: (a) The parabolic growth constants or diffusion coefficients of the thin film specimens are approximately twice as large as those of the single crystal specimens. The difference decreases with increasing temperature. (b) The activation energy for the E(Cu3 Sn) growth is 20% larger for the single crystal specimen. (c) The De°fvalue for the single crystal experiments is higher. (d) The lateral grain size of the e phase is larger for the single crystal experiments. The difference between the thin film and single crystal experiments decreases with increasing temperature. Observation (a) is explained as follows: Because of the small grain size, the effective interdiffusion coefficient, Oef, of a solid is the weighted average of the bulk diffusion coefficient, D, and the grain boundary diffusion coefficient, Db Def ---- D + a ( 2 6 /d)Db
(8)
where d is the average grain size, 6 is the thickness of the grain boundary and a is a shape constant (a ~ 1). As the grain size decreases, Oef increases because Db >>D. There may be several different reasons for observation (b), as Def is determined by some temperature dependent variables. Both the bulk and grain boundary diffusion coefficients, as well as the grain size, are temperature dependent and increase with rising temperature. Therefore the activation energy, Qer, is not a true constant (Fig. 10), although it can be considered as constant in a limited temperature range. If the temperature dependence of the grain size is neglected, then one explanation for observation (b) is that the activation energy of the bulk diffusion is higher than that of the grain boundary diffusion. Because of the larger grain size the bulk diffusion is responsible for a higher fraction of interdiffusion in the single crystal specimens than in the thin film specimens, and, therefore, the activation energy, Qef, is higher in the single crystal specimens. If it is assumed that, because of the small grain size ( < 1 pm), the grain boundary diffusion dominates in both cases in the temperature range studied, as it does in face-centered cubic metals (D b ~ 105 D at 0.5 Tin), a second possible reason for observation (b) could be that the temperature dependences of the E grain size in the single crystal and thin film experiments are different (d). The temperature dependence of the grain size in the thin film specimens is greater than that in the single crystal specimens. Therefore, the activation energy of diffusion is lower for the thin film specimens because of the reciprocal dependence of the effective diffusion coefficient on the grain size (Fig. 10). This assumption also explains observation
(c). In the Ni-Sn system (Fig. 2), both NisSn2 and Ni 3Sn phases can occur. Because at the start of the
Fig. 11. Ni/8.6/~m Sn/Ni. (a) 1281 h at 573 K, and (b) 125 h at 673 K. heat treatment only the Ni3Sn 4 phase grows they form in a Ni/Ni3 Sn4 diffusion couple (as in the Cu-Sn system). The experiments were made at 513, 573 and 673 K. The adhesion of the Ni layer on the SiO2 substrate was poor, so that they delaminated from the substrate after longer annealing times. At 513 and 573 K only the Ni3 Sn2 phase appeared in the thin film specimens. It grows as a perfect layer [Fig. 1l(a)]. The Ni3Sn2 grains in the layer are clearly visible in polarized light. They grow as thin, long, columnar grains which reach through the entire layer. The grain size is smaller than 1/~m. After annealing for 1000 h at 513 K and 86h at 573 K, during which Ni3Sn 2 layers grow to 1/~m in thickness, cracks in the Ni/Ni3 Sn2 interface appeared. The following reasons may be responsible for this: (a) A poor adhesion of the Ni and Ni3Sn 2 layers to each other because of a high specific 0 6
....
/
o. 3
/
:y.
/I
,,o:?:j o
.... 0
i
?2n 2
log
3
t (r~)
Fig. 12. Growth of the Ni3Sn 2 phase.
q
336
BADER et al.: FORMATION OF INTERMETALLIC COMPOUNDS such as vacancies and dislocations, form in the lattice. Those defects disappear with increasing annealing time, and the bulk diffusion coefficient decreases. (c) Possibly Kirkendall voids make it more difficult for the diffusant to pass through one of the interfaces.
o
::1_
.0
E~
o-0.3 /
/ -0,6
I
.
673 K
~ Ni/Ni3Sn 4 d i f f u s i o n m
i
i
0
To check the possibility of reason (a) the grain coarsening of the Ni 3Sn 2 phase was examined. At 573 and 673 K, the grain size was large enough for the average lateral grain size, d, to be measured. The data were fitted with the equation
10g t (h) Fig. 13. Growth of the Ni3Sn phase at 673 K. interfacial energy of the Ni/Ni3Sn 2 interphase boundary. (b) When cooling down after the heat treatment, the forces on the interface caused by the difference in thermal expansion coefficient of the Ni and Ni 3Sn 2 layers become too high and cracks develop. (c) It is known that in the crystal lattices of the Ni3 Sn2 and Ni3 Sn phases Ni diffuses very much faster than Sn [11]. This can cause the formation of Kirkendall voids at the Ni/Ni3Sn2 interface. Such voids reduce the contact and make crack formation possible during annealing or cooling down.
d = do" t s
(9)
by linear regression. The following values for the grain coarsening constant and the grain coarsening exponent were determined: do = 0.30 and 0.61/tm/h" and s = 0.15 and 0.19 for 573 and 673 K, respectively. At 573 K the average grain size increases from 0.5 to 1.0#m within 2000h, and at 673 K from 0.7 # m to 1.3/~m within the first 50 h. If grain boundary diffusion is assumed to be dominant, equation (8) can be approximated by (8a)
Der ~ 2a ~Db/d.
The following equation describes the growth of an IM (10)
d A x / d t ~ 2 ~D b A c / d A x
The morphology supports point (b) being responsible for the cracks. At 513 K the Ni3Sn 2 layer is 1.6 p m thick after 5000 h and at 573 K 3.7/~m after 2000 h (Fig. 12). At 673 K the Ni3Sn phase appears beside the Ni3 Sn2 phase. Also the Ni3 Sn phase grows as a plane layer with thin columnar grains. The grain size is only half the size of that of the Ni 3Sn 2 phase. Cracking occurs neither at the Ni/Ni3 Sn interface nor at the Ni3Sn/Ni3S % interface. Thus, there are grounds for the assumption that the cracks at 513 and 573 K are not caused by Kirkendail voids. The Ni3Sn 2 layer grows to 3 # m within 50 h (Fig. 12) and the Ni3Sn layer has a thickness of 1.9/~m at this time (Fig. 13). After 14h at 673 K the Ni3Sn 4 p h a s e h a s completely changed into the Ni 3Sn: and Ni 3Sn phase in the thin film specimens with 2 x 2.6 p m Sn. Both Ni3 Sn2 layers meet in the middle of the interconnection zone [Fig. ll(b)]. After additional 90h the Ni3Sn 2 phase has completely transformed into the Ni3Sn phase. To determine the growth law, the average layer thickness was measured and fitted with equation (2). The determined growth constants are listed in Table 3. In no case is the parabolic growth law exactly valid. There are several possible reasons for this: (a) Because of the small grain size, grain boundary diffusion dominates. The grain growth causes the effective diffusion coefficient to be time dependent and to decrease in the course of time. (b) At the start of heat treatment when the growth velocity of the IMs is very high, many lattice defects,
where Ac is the homogeneity range of the IM. If equation (9) is combined with equation (10) and the result is integrated, the following expression for Ax results Ax
f 4 ~ D b Ac'~ 0"5 0.5(1 --S) ~t0-0--~) t
or A x ~ t °'5`1-')
(11)
A comparison of equation (11) with equation (2) shows that n = 0.5(1 - s).
(12)
Placing the values of s determined in equation (12) results in the calculated growth exponents listed in Table 3. At 673 K, the measured and calculated values correspond very well. Also at 573 K there is a correspondence within the error range (+20%). This confirms the assumption that grain boundary diffusion dominates and that the deviation from the parabolic growth law to smaller n values is mainly due to the grain coarsening. Table 3. Growth constants (ko) and growth exponents (n) of the Ni3Sn2 and Ni3Sn phases and grain growth exponents (s) of the Ni3Sn2 phase Phase T (K) ko (/zm/min") n Ni3S% 513 0.044 0.43 573 0.226 0.36 673 0.669 0.40 Ni3Sn 673 0.309 0.45 Phase Ni3Sn2
T (K) 573 673
s 0.15 0.19
n (calc.) 0.42 0.41
n (exp.) 0.36 0.40
BADER et al.: FORMATION OF INTERMETALLIC COMPOUNDS The single crystal experiments carried out in the Ni-Sn system at 513 and 673 K were not significantly different from the thin film experiments with respect to morphology and growth velocity. Surprisingly, the Ni 3Sn phase appeared along with the Ni 3Sn 2 phase at 573 K, in contrast to the thin film experiments. Further studies showed that the Ni3Sn phase has difficulties nucleating at low temperatures, which can be reduced by contaminants at the Ni/Sn interface. Because the single crystal specimens were produced by evaporating Sn onto the Ni surface, the Ni surfaces came into contact with air before the Sn was evaporated, whereas the Ni and Sn layers of the thin film specimens were subsequently sputtered without opening the sputtering chamber. During growth of the Cu- and Ni-rich IMs faults can appear in the interconnection zone mainly in the shape of cracks and break-outs (caused by the preparation) at the link where the two layers grow together. This is probably due to a weakness of the IM/IM grain boundaries along the link. IMs often show an appreciable brittleness of the grain boundaries. This is due to the fact that the bonds in IMs are partly covalent and, therefore, oriented. At the grain boundaries the propability is low that the oriented bonds of two neighbouring grains match exactly. The grain boundaries of the long, thin, columnar grains are freely grown. This means that the bonds between two grains are relatively good. The grain boundaries of the joint where the two layers grow together, are not freely grown. Their position and orientation are determined by the growing together of the two layers. This is the reason that the grain boundary energy and brittleness are higher than the average there. The link of the two layers is a possible weak point of the interconnection. Cracks and break-outs can be reduced by applying a sufficient high pressure on the specimens while annealing.
4. CONCLUSIONS The experiments in the Cu-Sn and Ni-Sn systems led to the following results: 1. In the presence of liquid Sn, the IMs grow in morphologies which allow quick diffusion of Cu and Ni through the Sn(L). In these cases the growth of the IMs is proportional to t n with 0.2 < n < 0.3. 2. After the Sn(L) is changed into the Sn-rich IMs, the Cu- or Ni-rich IMs grow further in the shape of planar layers with small columnar grain size. In doing so, they deviate more or less from the
337
parabolic growth law. The E(Cu3Sn) and Ni3Sn phases obey the parabolic growth law relatively well (0.45 < n < 0.49). The Ni3 Sn2 phase, however, deviates appreciably from the parabolic growth law (0.36 < n < 0.43). This is due to the fact that the IMs have a very small grain size and grain boundary diffusion dominates. Because of the grain coarsening, the effective diffusion coefficient decreases with time. 3. The Cu3 Sn phase grows in the thin film specimens with double the velocity of its growth in the single crystal specimens. This is probably due to the larger grain size of the Cu3Sn phase developing in the single crystal specimens. 4. Nucleation is difficult for the Ni3Sn phase at temperatures below 623 K. Nucleation can be enhanced by contamination at the Ni/Sn interface. 5. The Ni/Ni 3Sn 2 interface, which exists when the Ni 3Sn phase does not nucleate, is mechanically weak. The evidence for this is the presence of cracks at the interface which form while cooling down after heat treatment. 6. Faults in the interconnection zone which reduce the reliability of the interconnection can form during annealing in the shape of pores caused by the volume contraction. During the growth of the Cu and Ni rich IMs cracks can form at the surface of contact between the two layers. By use of sufficient high pressures on the interconnection zone the formation of cracks can be suppressed. REFERENCES
1. H. Hieber, A. Swiderski, S. Bader and W. Gust, Proc. 7th European Hybrid Microelectronics Conf., Hamburg, Chapter 1.4. (1989). 2. S. Bader, W. Gust and H. Hieber, DVS-Berichte 122, 16 (1989). 3. S. Bader, W. Gust and H. Hieber, DVS-Berichte 129, 227 (1990). 4. H. Opielka, Prakt. Metallogr. 20, 388 (1983). 5. H. E. Bfihler and H. P. Hougardy, Atlas der Interferenzschichten-Metallographie. Deutsche Gesellschaft ffir Metallkunde, Oberursel (1979). 6. W. Villars and C. D. Calvert, Pearson's Handbook of Crystallographic Data for Intermetallic Phases. Am, Soc. Metals, Metals Park, Ohio (1985). 7. S. Tosima, S. Harada and E. O. Johnson, RCA (Radio Corporation of America) Review 45, 90 (March 1984). 8. C. Wagner, Acta metall. 17, 99 (1969). 9. T. B. Massalski et al. (editors), Binary Alloy Phase Diagrams. ASM Int., Materials Park, Ohio (1990). 10. C. A. Ma and R. A. Swalin, Acta metall. 8, 388 (1960). ll. J. A. van Beek, S. S. Stolk and F. J. J. van Loo, Z. Metallk. 73, 439 0982).
~
Pergamon
Copyright © 1994ElsevierScienceLtd Printed in Great Britain.All rights reserved 0956-7151/94$7.00+ 0.00
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RAPID FORMATION OF INTERMETALLIC C O M P O U N D S BY I N T E R D I F F U S I O N IN THE Cu-Sn A N D Ni-Sn SYSTEMS S. BADER1, W. GUST1 and H. HIEBER2 JMax-Planck-Institut ffir Metallforschung and Institut f/ir Metallkunde, Seestr. 75, D-70174 Stuttgart and 2Centrum fiir Mikroverbindungstechnik in der Elektronik, Forschung und Entwicklung GmbH, Ilsahl 5, D-24503 Neumiinster, Germany (Received 9 March 1994)
Abstract--Fundamental investigations were carried out to examine a new interconnection technology which is based on the rapid formation of intermetallic compounds composed of a high melting component (e.g. Cu or Ni) and a low melting component (e.g. Sn) between two layers of the high melting component at temperatures just above the melting point of Sn. This reaction is known as isothermal solidification. The growth of the intermetallic compounds of the Cu-Sn and Ni-Sn systems were studied in thin films at temperatures from 513 to 673 K. At the beginning of interdiffusion, the Sn-rich intermetallic compounds r/ (Cu6Sn5) and Ni3Sn4 grow fastest with a non-parabolic time dependence. The Cu3Sn phase grows parabolically, however, not the Ni3Sn2 phase. The Ni3Sn phase does not nucleate below 623 K in specimens with clean Ni/Sn interfaces. For the first time, the influence of the small grain size of the Cu or Ni thin filmswere studied by performing similar experimentswith Cu or Ni singlecrystals. In the Cu-Sn system the interdiffusion coefficients for Cu3Sn obtained from the thin film experiments are twice those obtained from the single crystal experiments. In the Ni-Sn system there are no differences between thin film and single crystal results.
1. INTRODUCTION The formation and growth of intermetallics (IMs) in thin C u - S n - C u and N i - S n - N i films have been studied. The reason for these investigations was the examination of a new thin film contact technology [1-3]. This technology is based on the rapid formation of IMs composed of a high melting component (HC), such as Cu and Ni, and a low melting component (LC), such as Sn, at temperatures just above the melting point of the LC between two layers of the HC. Figure 1 shows this schematically. The bond-layer grows by liquid-solid interdiffusion at the thin film interfaces and may consist of one or more IMs. After consumption of the liquid LC the interconnection is stable up to the melting temperature, Tin, of the IM with the lowest melting point existing in the bond-layer. Figure 2 shows this for the Ni-Sn system. By variation of the film thicknesses of the HC and LC and the time of the heat treatment, the kind and proportionate amount of IMs can be influenced. In this way small and heat-resistant interconnections can be produced. In contrast to the soft-solder technology the maximal use temperature is here appreciably higher than the bonding temperature, Tb, at which the bond is produced Tm> Tb + 500 K.
(1)
This technology can therefore be applied to the packaging of devices which exhibit high heat dissi-
pation or which are exposed to high temperatures or changes in ambient temperature. The feasibility of this technology depends on the growth velocity, the aging behavior and the morphology of the IMs in the bonding layer. Initial experiments showed that all possible systems HC-LC exhibited high growth velocities for the IMs. The Cu-Sn and Ni-Sn systems were choosen for further studies because they have only a few IMs and the compositions and structures of the IMs are all well-known.
2. EXPERIMENTALPROCEDURE The specimens were made by sequential sputtering of Cu or Ni and Sn layers on Si wafers (0.35 mm thick). The specimens had dimensions of 4 x 4 × 0.35 mm 3 and were made by cleaving from a appropriately scribed Si wafer. Because of the low adhesion of Cu on Si a Ni-Cr-AI adhesion layer was used. To avoid any influence of the adhesion layer on the growth of the IMs, the Cu layers were made thicker. To produce the films, a 2400 type Perkin Elmer sputter depositer was used. The base pressure before sputtering was 5.3 x 10-4 Pa and the Ar pressure while sputtering was 1.3 Pa. The following combinations of film thicknesses were produced: Cu(7/~m)/Sn(2.6/~m), Cu(7 pm)/Sn(l.6/~m), Ni(4.4/~m)/Sn(4.3/tm) and Ni(4.1 #m)/Sn(2.6/~m). During sputtering of the Sn layers onto Cu, an
329
330
BADER et aL: FORMATION OF INTERMETALLIC COMPOUNDS
approx. 0.3 # m thick layer of IMs grew. In the Ni-Sn system there was no growth of IMs while sputtering. Sputtered layers of the HC have a very small grain size. To study the influence of this small grain size, similar experiments using Cu and Ni single crystals were done. For these experiments Sn layers comparable to the Sn layers of the sputtered specimens were evaporated onto polished single crystal specimens having the dimensions 4 x 4 x 1 mm 3. Two pieces were put together as shown in Fig. 1 and placed in a specimen holder, which exerted a small but uniform pressure on the specimen through a central screw. This "two piece" arrangement caused the heating up time to be longer because of the heat capacity of the specimen holder, however, it prevented the formation of Sn droplets while melting, enlarged the useable interface length for measuring the width of the ICs and was very similar to a later application. Heat treatments up to 20 min were carried out in a big (186 cm 3) Pb-Sn solder bath. In this case, the specimens were wrapped in A1 foil to prevent reactions with the Pb-Sn solder. The effective heating up time was only 2-3 s and could be neglected in the analysis. Longer annealings were carried out in an oil bath or in tube furnaces. Here the specimens with specimen holders were sealed in evacuated glass tubes backfilled with 200 mbar Ar to protect them against oil or corrosion. Most of the observations were made using an optical microscope (LM). Only in special cases was the scanning electron microscope (SEM) used. The annealed specimens were mounted in epoxy and ground perpendicular to the metal and IM layers with 600, 1200 and 4000 grit abrasive paper. Polishing followed with 3 and 1 # m diamond paste and colloidal SiO 2 suspension [4]. If Sn was still visible after
Before reaction
HC
HC
Sub
Sub
After reaction
SubI
t F--
5
0
0
°°°I/ 3
0 N1
~
T2
\
......:::1 0
20
0
40 at.%
~
60 Sn
80
100 Sn
Fig. 2. Ni-Sn phase diagram [9]. If there is Ni3Sn4 phase in the bond area, the bond is stable up to T~. If there is only Ni3Sn2phase in the bond area, the bond is stable up to T2. L = liquid phase, Tb = bonding temperature. the heat treatment, the IM/Sn interface could be made more visible by selectively etching away the Sn with an aqueous solution of 35 g/l ortho-nitrophenol and 50 g/1 NaOH. The ability to distinguish between the other phases could not be improved by chemical and electrochemical methods. The interference layer method [5] was the solution to this problem. Here a thin oxide film is reactively sputtered on the surface of the specimen to be used as a reflection interference filter. In this way, the bright~lark and color contrasts are improved. For the Cu-Sn specimens a Pt anode was used, and for the Ni-Sn specimens a Pb anode. For the optical microscope observations, a Zeiss Axiophot-type microscope with an oil differential interference contrast objective and a magnification of 1000:1 was used. The width of the IM layers were measured on the screen of a video system connected to the microscope at a total magnification of 5000: 1. Because some of the IM interfaces grew unevenly and did not remain planar the average thickness of an IM layer was determined by measuring the layer thickness at 50 equally spaced points and calculating the average thickness. The measurements are known to within +0.2 #m. The standard deviation for different IM layers was between 0.05 and 0.15 #m. The identification of the IMs were carried out by X-ray diffraction measurements. A Guinier camera with asymmetric transmission arrangement was used. The multilayer thin films were stripped off from their substrates and examined directly. For identification, the measured scans were compared with calculated scans for the different ICs. The data for the calculations were obtained from the literature [6]. In all cases, the IMs could be identified as the stable IMs of the phase diagrams.
HC
IM
Sub
3. RESULTS AND DISCUSSION
3.1. Growth of Sn-rich intermetallic compounds in the presence of liquid Sn C
Fig. 1. Structure of the bond area before and after heat treatment. HC, LC = high and low melting component, Sub = substrate, IM = intermetallic compound.
The IMs grow by liquid-solid interdiffusion. First, the changes in the border phases will be considered. The saturation of the Sn(L) is very quick, because the saturation concentrations of the HC in Sn(L) are low
BADER et al.:
FORMATION OF INTERMETALLIC COMPOUNDS
in the temperature range studied (513-603 K) and the diffusion in the liquid is very fast. An estimation and measurements by Tosima e t al. [7] showed that the saturation time is between 1 and 3 s. The saturation concentration of Sn in the HC is very low as well, and the measurements showed that the decrease in thickness of the HC at the HC/IM interface for the IM formation is higher than the penetration depth of Sn in the HC. So the border phases are the pure HC and saturated Sn(L). The beginning of the Cu-Sn interdiffusion was studied at 513 and 573 K using the specimens with a 2 × 2.6/~m thick Sn layer. The r/ and E phases form simultaneously [Fig. 3(a)]. The growth velocity of the r/phase is very fast compared to that of the E phase. While the E phase grows in the form of a continuous and wavy layer, the r/phase grows in a hemispherical shape [Fig. 3(a, b)]. The hemispheres were identified
as single crystals using polarized light. Between them are often deep grooves down to the Cu. At these places the thickness of the Cu- and E-layer are minimums. Therefore, there are grounds for the supposition that these grooves are important diffusion paths for Cu and Sn in addition to bulk diffusion through the ~/ phase. The Cu enters the Sn(L) at the base of the grooves and moves quickly because of the high diffusivity and convection to the top of the hemispheres. With time, the hemispheres grow perpendicular to the Cu surface and by a process comparable to grain growth parallel to it, while the hemispherical shape is preserved. The growth exponent of the grain growth is similar to that of particle coarsening and is predicted by the Lifshitz-Wagner law. The grain size at higher temperature is always larger than that at lower temperature. After 8-10min at 513 K, the Sn is changed completely into the IMs [Fig. 3(c)]. The ~/ crystals extend across the interconnection area. The grain boundaries lie perpendicular to the layers. A faultless and voidless interconnection zone has grown between the Cu layers. The ~/ crystallites extend through the entire zone from one Cu layer to the other. Defects in the interconnection zone can form in the shape of pores. They form for instance if the pressure on the specimen is too low during annealing, because of the volume contraction, which occurs during formation of the IMs and which is 7% for the Cu-Sn system. The pores often form in the shape of long thin gaps (fissures) between the r/ crystallites, when there is insufficient Sn for the crystals to grow together. The shape of the hemispheres made it difficult to measure the thickness, Ax, of the r/ phase and increased the error in these measurements. The average thickness was determined as described above. The annealing times, t, varied from 5 s to 5 min. Until the r/ phase reached a thickness of 2.0-2.5 #m, the distance between the r/ layers on both sides were large enough so they did not influence each other. The diffusion-controlled growth of an IM normally obeys a parabolic growth law [8] A x = ko " t ~
Fig. 3. 513 K, Cu/5.2/~m Sn/Cu (a-c) and Cu/3.2/~m Sn/Cu (d) after I min (a), 30s (b), 12min (c) and 20min (d). Sn selectively etched away (a, b).
331
or
l o g A x = l o g ko ÷ n l o g t
(2)
where n equals 0.5. Table 1 shows the measured widths of the ~/and E layers. The data from the 513 K experiments are the average of 2-4 measurements. In Fig. 4, the data for the ~/ phase are plotted in a log Ax-log t diagram. The n and k0 values computed from the curves in this figure by linear regression are the following: n = 0.21 and 0.25 and k 0 = 1.906 and 1.800/tm/min n for 513 and 573 K, respectively. The following two observations were noted: (a) the growth of the t/ phase does not obey a parabolic growth law at the beginning of interdiffusion; (b) at short annealing times, the average widths of the r/ layers for the 513 K experiments are larger than those for the 573 K experiments. The deviation from the parabolic growth law
332
BADER et al.:
FORMATION OF INTERMETALLIC COMPOUNDS
Table 1. Average thickness of the r/- and E-phase layers in the experiments where liquid Sn was still present Layer width (,am) Experiment t at 513 K (min) Ax, Ax~ Ax'~ 1-3 0.083 0.15 1.15 1.25 4-6 0.25 0.20 1.40 1.59 7-10 0.50 0.26 1.68 1.85 11-13 1.00 0.30 1.85 2.06 14-15 2.50 0.40 2.35 2.63
t
(min) 0.17 0.18 0.53 0.50 1.00 1.00 2.00 2.00
Ax, 0.25 0.25 0.58 0.65 0.75 0.75 0.80 0.83
Ax~ 1.27 1.18 1.62 1.47 1.82 1.74 2.12 2.17
Ax~ = Ax~ + Axe,.
(3)
In this equation Axe, is part of the q-layer thickness, which was used to form the e phase. The following reaction is responsible for the formation of the e phase 9 Cu + C u 6 Sn5 ~ 5 Cu3 Sn.
Layer width (#m) Experiment at 573 K 1 2 3 4 5 6 7 8
F o r this reason an q-layer thickness Ax~ was calculated using the equation
Ax~ 1.44 1.35 2.02 1.92 2.34 2.26 2.67 2.74
Thus the following expression for Ax,, results Ax,, = 11 V, Ax,/20 V,
(4)
where V, and V, are the molar volumes of the q and c phases, respectively. If the reduction of the q phase is taken into account by using Ax~ instead of Ax~ in Fig. 4, the point of intersection of the two lines lies
cannot be attributed only to the fact that the IM layer formed during deposition was ignored. A n estimation showed this not to be the case. The reason for the deviation is probably that a large part of transport of the H C and Sn to the phase boundaries, where the formation of the q phase takes place, runs through the grooves between q crystallites. In the course of time, these grooves get narrower and their density is reduced by a process comparable to grain growth or Ostwald ripening. This means that the diffusion slows down not only by the increase in diffusion path length and reduction in concentration gradient, but also by the reduction of transport capacity of the grooves. This is manifested in n values less than 0.5. Our explanation for the second observation is that in addition to the q phase, the e phase is also growing. As mentioned above, the q phase also grows by diffusion of copper through the liquid Sn. The activation energy for the Cu diffusion in Sn(L) is only 17.5 kJ/mol [10]. The q phase is both produced by reaction of E phase with Sn and is reduced to E phase by reaction of q phase with Cu. The activation energy for interdiffusion in the E phase is 65 kJ/mol. This means that the ratio of q growth rate is smaller at higher temperatures.
0.%
0.4 -E ._~0.3 t~0.2 0
0.1 0.0
- i .2 ]og t (min)
Fig. 4. Growth of the r/(Cu6Sns) phase at 513 and 573 K.
Fig. 5. Ni/8.6 #m Sn/Ni at 513 K after 165 s (a), 60 s (b) and 90 min (c). Sn selectively etched away (a, b); section parallel to intermetallic layers (b).
BADER et al.:
-0.1
-1.0
FORMATION OF INTERMETALLIC COMPOUNDS
5"13 K
-0.5
0.0
0.5
.0
log t (rain) Fig. " .~erease of the Ni film thickness during the reaction to form Ni3Sn4 as a function of time at 513 K.
at a time of 20 s. In addition, the saturation concentration of Cu in Sn(L) is higher at higher temperatures; therefore, more time is necessary to saturate the Sn. The lateral size of the r/ hemispheres is larger at higher temperatures, thus decreasing the total amount of fast diffusion path (grooves). Therefore, the growth rate of the r/ phase is smaller at higher temperatures in the initial stages of annealing. The initial stages of growth in the Ni-Sn system were studied by using specimens with a 2 x 4.3 p m thick Sn layer. Figure 5(a, b) shows SEM photographs perpendicular and parallel, respectively, to the layers of specimens with short annealing times. At the start of annealing, only one of the three IMs of the Ni-Sn system (Fig. 2) appears. This initial IM was identified as the Ni3Sn4 phase and was found to exhibit three different morphologies simultaneously: as a fine-grained, planar layer at the Ni interface; long, thin, idiomorphic whiskers; and large, polygonal, idiomorphic crystals. After an annealing time of only 7 s, a high density of Ni3Sn 4 whiskers is present. The large crystals are not yet seen. The density and thickness of the whiskers reach their maximum after 2 min, after which they decrease. After 30 s, the first large crystals appear and their number increases rapidly from the first to fifth minute of annealing. Obviously a recrystallization of the whiskers into the large crystals takes place. The driving force for this recrystallization is the reduction of surface area and energy of the Ni3 Sn4 phase. Meanwhile the fine-grained layer grows obeying parabolic growth law. After 60 min, the Sn is completely changed into Ni3Sn 4. The entire layer now consists of two equally thick ( ~ 1 #m) finegrained layers at the Ni interfaces and a coarse grained area in the middle [Fig. 5(c)]. When the pressure on the specimen was too low during annealing, there is a band of small pores between these two areas. The quick growth and the idiomorphic shape of the whiskers and big crystals indicate that the N i 3 S n 4 phase initially forms almost exclusively by the diffusion of Ni through the Sn(L) and crystallization from the liquid phase. The low solubility of Ni in Sn(L) and the high density of whiskers preclude AM 43/I--V
333
whiskers from forming during cooling. Obviously the NiaSn4 phase has a special lattice direction with a high crystallization velocity. Grains with the appropriate orientation grow out of the fine-grained layer and form the whiskers. The shape of the Ni 3Sn 4 phase made it impossible to measure the average layer width directly. The relative amount of Ni3Sn 4 phase formed was determined by measuring the decrease of the Ni layer thickness, ANi. Because the thickness of the Ni layer was not uniform across the Si wafer, the thickness of each specimen was measured before and after the heat treatment. Therefore the error of ANi was relatively high. Figure 6 shows the log Ax-log t plot of the measured data. When the data were fit with a straight line by linear regression the following k and n values were obtained: n = 0.28, k = 0.453 #m/min °'28. As was the case for the Cu-Sn, the growth exponent is smaller than 0.5, and the reason for this behavior is the same as that for Cu-Sn. In the course of time the transport capacity of the diffusion paths through the Sn(L) decreases because of the increasing density of whiskers and large crystals. The morphology in both systems is shaped by the fact that part of the Cu and Sn diffusion goes through the Sn(L). Because of the higher interface area, these morphologies are energetically less favorable; however, they result in an increase in the temporal reduction of free energy. Probably the morphology that is formed is mainly determined by the relation of the effective interdiffusion coefficients [see equation (8)] or transport capacities of the IMs and the Sn(L). If the effective interdiffusion coefficients, Def, of the IMs are low, then the transport through the Sn(L) predominates. This results in the rugged or open morphology and the idiomorphic shapes of the IC, which form by crystallization from the liquid phase. The melting point of Ni3 Sn4 is 1087 K, which means that, at 513 K, Oef of the Ni3Sn 4 phase is probably lower than D~f of the r/ phase which has a melting point of 688 K. Therefore the transport of the HC through the Sn(L) is probably larger in the case of Ni-Sn than in the case of Cu-Sn. This is visible in the
0 . 4 F
0.5
.......
T ........... ~
0.0
0.5 ]og
--7
1.0
]5
0
t (rnin)
Fig. 7. Growth of the E(Cu 3Sn) phase in the temperature range from 513 to 603 K.
BADER et al.: FORMATION OF INTERMETALLIC COMPOUNDS
334
600
500 K
-13
© E
Fig. 9. Cu/3.2 #m Sn/Cu, 30 min at 543 K. -15 15
1'8 i i04/T (I/K) Fig. 8. Arrhenius plot of Def of the thin film and single crystal experiments. morphology. Other factors which influence the morphology are the specific interface energy of the IM/Sn(L) boundary and the anisotropy of the effective interdiffusion and of the crystallization.
3.2. Growth o f Cu- or Ni-rich intermetallic compounds In the Cu-Sn system the growth of the e(Cu3Sn) phase was studied in the temperature range from 513 to 603 K using the specimens with 2 x 1.6 # m thick Sn layers. The Sn changes to the r/phase within a few seconds, and the slower growing E phase grows as it does in a Cu/r/(Cu6 Sn5) diffusion couple. The E phase grows in a planar shape [Fig. 3(d)]. At short annealing times and low temperatures no grain structure is visible. For longer heat treatments and higher temperatures a very fine grain structure of thin, columnar grains, which reach through the entire layer, is visible under polarized light. The lateral grain size (parallel to the layers) is less than 1 p m and was too small to be measured. The measured data for the layer thickness are plotted in Fig. 7. The E phase grows to a layer thickness of 2 # m within 40 rain at 513 K and within 4 min at 603 K. The data were fitted with equation (2), and the n values were found between 0.45 and 0.49. Therefore the parabolic growth law was assumed to be valid. The parabolic growth constants, kp
kp = Ax2/2t
To compute the effective interdiffusion coefficients, equations derived by Wagner [8] were used. For interdiffusion in the E phase in a Cu/r/(Cu6Sns) diffusion couple the following equation is valid Def = ( c , - Ccu)(C~ - _c,)kp
(6)
Ac,(c, - Cc.)
where c is the mole fraction of Sn, c, is the average composition of the E phase, Cc. is the saturation concentration of Sn in Cu (Cc. = 0), c, is the Sn concentration of the r/phase in equilibrium with the E phase and Ac, is the homogeneity range of the E phase. The concentrations are taken from the phase diagram [9]. The Def values are plotted in Fig. 8 and fitted with an Arrhenius equation
Oet= O°f exp(-- Q,f/RT)
(7)
where R is the gas constant and T is the absolute temperature. The determined values of the preexponential factor, De°, and the activation energy, Qef, are listed in Table 2. Figure 9 shows a specimen annealed for 30 min at 543 K. Both e layers have grown together. The interconnection zone is widely free of pores. To study the influence of the small grain size of the thin films, similar experiments with Cu single crystals were made. The morphology of the E layer was the same. The growth of the E phase obeys the parabolic growth law too (0.48 < n < 0.49). The calculated D~f values are plotted in Fig. 8. The following differences
(5)
determined are listed in Table 2.
Table 2. Parabolicgrowth constant (kp) and Arrheniusparameters (Deft Qef) for the Ephase kp (10 lSm2/s) T (K) Thin film Singlecrystal 513 0.92 0.38 543 2.00 0.94 573 4.64 2.85 603 9.30 -Specimens Thin film Singlecrystal
D~r(10 8m2/s) 5.3 80.8
Qef(kJ/mol) 66.1 81.6
lIT Fig. 10. Schematic description of the temperature dependence of De of the E(Cu3Sn) phase. A = temperature range studied; D0f Xr, Dsc = D,f of the thin film and single crystal experiments; D, Db = bulk and grain boundary diffusion coefficients.
BADER et al.: FORMATION OF INTERMETALLIC COMPOUNDS
335
between the thin film and single crystal experiments were observed: (a) The parabolic growth constants or diffusion coefficients of the thin film specimens are approximately twice as large as those of the single crystal specimens. The difference decreases with increasing temperature. (b) The activation energy for the E(Cu3 Sn) growth is 20% larger for the single crystal specimen. (c) The De°fvalue for the single crystal experiments is higher. (d) The lateral grain size of the e phase is larger for the single crystal experiments. The difference between the thin film and single crystal experiments decreases with increasing temperature. Observation (a) is explained as follows: Because of the small grain size, the effective interdiffusion coefficient, Oef, of a solid is the weighted average of the bulk diffusion coefficient, D, and the grain boundary diffusion coefficient, Db Def ---- D + a ( 2 6 /d)Db
(8)
where d is the average grain size, 6 is the thickness of the grain boundary and a is a shape constant (a ~ 1). As the grain size decreases, Oef increases because Db >>D. There may be several different reasons for observation (b), as Def is determined by some temperature dependent variables. Both the bulk and grain boundary diffusion coefficients, as well as the grain size, are temperature dependent and increase with rising temperature. Therefore the activation energy, Qer, is not a true constant (Fig. 10), although it can be considered as constant in a limited temperature range. If the temperature dependence of the grain size is neglected, then one explanation for observation (b) is that the activation energy of the bulk diffusion is higher than that of the grain boundary diffusion. Because of the larger grain size the bulk diffusion is responsible for a higher fraction of interdiffusion in the single crystal specimens than in the thin film specimens, and, therefore, the activation energy, Qef, is higher in the single crystal specimens. If it is assumed that, because of the small grain size ( < 1 pm), the grain boundary diffusion dominates in both cases in the temperature range studied, as it does in face-centered cubic metals (D b ~ 105 D at 0.5 Tin), a second possible reason for observation (b) could be that the temperature dependences of the E grain size in the single crystal and thin film experiments are different (d). The temperature dependence of the grain size in the thin film specimens is greater than that in the single crystal specimens. Therefore, the activation energy of diffusion is lower for the thin film specimens because of the reciprocal dependence of the effective diffusion coefficient on the grain size (Fig. 10). This assumption also explains observation
(c). In the Ni-Sn system (Fig. 2), both NisSn2 and Ni 3Sn phases can occur. Because at the start of the
Fig. 11. Ni/8.6/~m Sn/Ni. (a) 1281 h at 573 K, and (b) 125 h at 673 K. heat treatment only the Ni3Sn 4 phase grows they form in a Ni/Ni3 Sn4 diffusion couple (as in the Cu-Sn system). The experiments were made at 513, 573 and 673 K. The adhesion of the Ni layer on the SiO2 substrate was poor, so that they delaminated from the substrate after longer annealing times. At 513 and 573 K only the Ni3 Sn2 phase appeared in the thin film specimens. It grows as a perfect layer [Fig. 1l(a)]. The Ni3Sn2 grains in the layer are clearly visible in polarized light. They grow as thin, long, columnar grains which reach through the entire layer. The grain size is smaller than 1/~m. After annealing for 1000 h at 513 K and 86h at 573 K, during which Ni3Sn 2 layers grow to 1/~m in thickness, cracks in the Ni/Ni3 Sn2 interface appeared. The following reasons may be responsible for this: (a) A poor adhesion of the Ni and Ni3Sn 2 layers to each other because of a high specific 0 6
....
/
o. 3
/
:y.
/I
,,o:?:j o
.... 0
i
?2n 2
log
3
t (r~)
Fig. 12. Growth of the Ni3Sn 2 phase.
q
336
BADER et al.: FORMATION OF INTERMETALLIC COMPOUNDS such as vacancies and dislocations, form in the lattice. Those defects disappear with increasing annealing time, and the bulk diffusion coefficient decreases. (c) Possibly Kirkendall voids make it more difficult for the diffusant to pass through one of the interfaces.
o
::1_
.0
E~
o-0.3 /
/ -0,6
I
.
673 K
~ Ni/Ni3Sn 4 d i f f u s i o n m
i
i
0
To check the possibility of reason (a) the grain coarsening of the Ni 3Sn 2 phase was examined. At 573 and 673 K, the grain size was large enough for the average lateral grain size, d, to be measured. The data were fitted with the equation
10g t (h) Fig. 13. Growth of the Ni3Sn phase at 673 K. interfacial energy of the Ni/Ni3Sn 2 interphase boundary. (b) When cooling down after the heat treatment, the forces on the interface caused by the difference in thermal expansion coefficient of the Ni and Ni 3Sn 2 layers become too high and cracks develop. (c) It is known that in the crystal lattices of the Ni3 Sn2 and Ni3 Sn phases Ni diffuses very much faster than Sn [11]. This can cause the formation of Kirkendall voids at the Ni/Ni3Sn2 interface. Such voids reduce the contact and make crack formation possible during annealing or cooling down.
d = do" t s
(9)
by linear regression. The following values for the grain coarsening constant and the grain coarsening exponent were determined: do = 0.30 and 0.61/tm/h" and s = 0.15 and 0.19 for 573 and 673 K, respectively. At 573 K the average grain size increases from 0.5 to 1.0#m within 2000h, and at 673 K from 0.7 # m to 1.3/~m within the first 50 h. If grain boundary diffusion is assumed to be dominant, equation (8) can be approximated by (8a)
Der ~ 2a ~Db/d.
The following equation describes the growth of an IM (10)
d A x / d t ~ 2 ~D b A c / d A x
The morphology supports point (b) being responsible for the cracks. At 513 K the Ni3Sn 2 layer is 1.6 p m thick after 5000 h and at 573 K 3.7/~m after 2000 h (Fig. 12). At 673 K the Ni3Sn phase appears beside the Ni3 Sn2 phase. Also the Ni3 Sn phase grows as a plane layer with thin columnar grains. The grain size is only half the size of that of the Ni 3Sn 2 phase. Cracking occurs neither at the Ni/Ni3 Sn interface nor at the Ni3Sn/Ni3S % interface. Thus, there are grounds for the assumption that the cracks at 513 and 573 K are not caused by Kirkendail voids. The Ni3Sn 2 layer grows to 3 # m within 50 h (Fig. 12) and the Ni3Sn layer has a thickness of 1.9/~m at this time (Fig. 13). After 14h at 673 K the Ni3Sn 4 p h a s e h a s completely changed into the Ni 3Sn: and Ni 3Sn phase in the thin film specimens with 2 x 2.6 p m Sn. Both Ni3 Sn2 layers meet in the middle of the interconnection zone [Fig. ll(b)]. After additional 90h the Ni3Sn 2 phase has completely transformed into the Ni3Sn phase. To determine the growth law, the average layer thickness was measured and fitted with equation (2). The determined growth constants are listed in Table 3. In no case is the parabolic growth law exactly valid. There are several possible reasons for this: (a) Because of the small grain size, grain boundary diffusion dominates. The grain growth causes the effective diffusion coefficient to be time dependent and to decrease in the course of time. (b) At the start of heat treatment when the growth velocity of the IMs is very high, many lattice defects,
where Ac is the homogeneity range of the IM. If equation (9) is combined with equation (10) and the result is integrated, the following expression for Ax results Ax
f 4 ~ D b Ac'~ 0"5 0.5(1 --S) ~t0-0--~) t
or A x ~ t °'5`1-')
(11)
A comparison of equation (11) with equation (2) shows that n = 0.5(1 - s).
(12)
Placing the values of s determined in equation (12) results in the calculated growth exponents listed in Table 3. At 673 K, the measured and calculated values correspond very well. Also at 573 K there is a correspondence within the error range (+20%). This confirms the assumption that grain boundary diffusion dominates and that the deviation from the parabolic growth law to smaller n values is mainly due to the grain coarsening. Table 3. Growth constants (ko) and growth exponents (n) of the Ni3Sn2 and Ni3Sn phases and grain growth exponents (s) of the Ni3Sn2 phase Phase T (K) ko (/zm/min") n Ni3S% 513 0.044 0.43 573 0.226 0.36 673 0.669 0.40 Ni3Sn 673 0.309 0.45 Phase Ni3Sn2
T (K) 573 673
s 0.15 0.19
n (calc.) 0.42 0.41
n (exp.) 0.36 0.40
BADER et al.: FORMATION OF INTERMETALLIC COMPOUNDS The single crystal experiments carried out in the Ni-Sn system at 513 and 673 K were not significantly different from the thin film experiments with respect to morphology and growth velocity. Surprisingly, the Ni 3Sn phase appeared along with the Ni 3Sn 2 phase at 573 K, in contrast to the thin film experiments. Further studies showed that the Ni3Sn phase has difficulties nucleating at low temperatures, which can be reduced by contaminants at the Ni/Sn interface. Because the single crystal specimens were produced by evaporating Sn onto the Ni surface, the Ni surfaces came into contact with air before the Sn was evaporated, whereas the Ni and Sn layers of the thin film specimens were subsequently sputtered without opening the sputtering chamber. During growth of the Cu- and Ni-rich IMs faults can appear in the interconnection zone mainly in the shape of cracks and break-outs (caused by the preparation) at the link where the two layers grow together. This is probably due to a weakness of the IM/IM grain boundaries along the link. IMs often show an appreciable brittleness of the grain boundaries. This is due to the fact that the bonds in IMs are partly covalent and, therefore, oriented. At the grain boundaries the propability is low that the oriented bonds of two neighbouring grains match exactly. The grain boundaries of the long, thin, columnar grains are freely grown. This means that the bonds between two grains are relatively good. The grain boundaries of the joint where the two layers grow together, are not freely grown. Their position and orientation are determined by the growing together of the two layers. This is the reason that the grain boundary energy and brittleness are higher than the average there. The link of the two layers is a possible weak point of the interconnection. Cracks and break-outs can be reduced by applying a sufficient high pressure on the specimens while annealing.
4. CONCLUSIONS The experiments in the Cu-Sn and Ni-Sn systems led to the following results: 1. In the presence of liquid Sn, the IMs grow in morphologies which allow quick diffusion of Cu and Ni through the Sn(L). In these cases the growth of the IMs is proportional to t n with 0.2 < n < 0.3. 2. After the Sn(L) is changed into the Sn-rich IMs, the Cu- or Ni-rich IMs grow further in the shape of planar layers with small columnar grain size. In doing so, they deviate more or less from the
337
parabolic growth law. The E(Cu3Sn) and Ni3Sn phases obey the parabolic growth law relatively well (0.45 < n < 0.49). The Ni3 Sn2 phase, however, deviates appreciably from the parabolic growth law (0.36 < n < 0.43). This is due to the fact that the IMs have a very small grain size and grain boundary diffusion dominates. Because of the grain coarsening, the effective diffusion coefficient decreases with time. 3. The Cu3 Sn phase grows in the thin film specimens with double the velocity of its growth in the single crystal specimens. This is probably due to the larger grain size of the Cu3Sn phase developing in the single crystal specimens. 4. Nucleation is difficult for the Ni3Sn phase at temperatures below 623 K. Nucleation can be enhanced by contamination at the Ni/Sn interface. 5. The Ni/Ni 3Sn 2 interface, which exists when the Ni 3Sn phase does not nucleate, is mechanically weak. The evidence for this is the presence of cracks at the interface which form while cooling down after heat treatment. 6. Faults in the interconnection zone which reduce the reliability of the interconnection can form during annealing in the shape of pores caused by the volume contraction. During the growth of the Cu and Ni rich IMs cracks can form at the surface of contact between the two layers. By use of sufficient high pressures on the interconnection zone the formation of cracks can be suppressed. REFERENCES
1. H. Hieber, A. Swiderski, S. Bader and W. Gust, Proc. 7th European Hybrid Microelectronics Conf., Hamburg, Chapter 1.4. (1989). 2. S. Bader, W. Gust and H. Hieber, DVS-Berichte 122, 16 (1989). 3. S. Bader, W. Gust and H. Hieber, DVS-Berichte 129, 227 (1990). 4. H. Opielka, Prakt. Metallogr. 20, 388 (1983). 5. H. E. Bfihler and H. P. Hougardy, Atlas der Interferenzschichten-Metallographie. Deutsche Gesellschaft ffir Metallkunde, Oberursel (1979). 6. W. Villars and C. D. Calvert, Pearson's Handbook of Crystallographic Data for Intermetallic Phases. Am, Soc. Metals, Metals Park, Ohio (1985). 7. S. Tosima, S. Harada and E. O. Johnson, RCA (Radio Corporation of America) Review 45, 90 (March 1984). 8. C. Wagner, Acta metall. 17, 99 (1969). 9. T. B. Massalski et al. (editors), Binary Alloy Phase Diagrams. ASM Int., Materials Park, Ohio (1990). 10. C. A. Ma and R. A. Swalin, Acta metall. 8, 388 (1960). ll. J. A. van Beek, S. S. Stolk and F. J. J. van Loo, Z. Metallk. 73, 439 0982).