Passivation of Cu via refractory metal nitridation in an ammonia ambient

$Page 1: Passivation of Cu via refractory metal nitridation in an ammonia ambient$
Thin Solid Films 262 (1995) 199-208

Passivation of Cu via refractory metal nitridation in an ammonia ambient

Daniel Adamsa, T.L. Alford”, N.D. Theodore”, S.W. Russell”, R.L. Spreitzer”“, J.W. Mayerh

“Department of Chemical, Bio and Materials Engineering, Arizona State University, Tempe, AZ X5287-6006. USA hCenter for Solid State Science, Arizona State University, Tempe, AZ 85287-I 704, USA

‘Intel Corporation. CH4-55. Chandler. AZ X.5226, USA

Abstract

Cu-(27at.s Ti) and Cu-(26at.% Cr) alloys codeposited on silicon dioxide substrates were isochronally annealed for 30 min at 400-700 “C in a flowing NH, ambient. In the Cu-Ti alloy, Ti segregates to the free surface to form a TIN.(O) layer and also to the alloy-SiOz interface to form a Ti,Si,/TiO, bilayer structure. Therefore the resulting structure is an almost completely dealloyed Cu layer located between a surface oxygen-rich Ti nitride and Ti-silicide/Ti-oxide bilayer interfacial structure. In the Cu-Cr alloy system, Cr seems to migrate only to the free surface to form a CrN, passivation layer. A 4.5 nm Al film was deposited after nitridation. whereupon a second anneal was performed to evaluate these nitride layers as diffusion barriers. The Cr-nitride diffusion barrier is stable up to 600 “C compared with Ti nitride that fails at 500 “C. The Cu-Cr nitrided samples also showed an overall lower sheet resistance.

Keywords: Ammonia; Diffusion; Titanium nitride; Titanium oxide

1. Introduction

Aluminum and its alloys are currently used as the primary contact and interconnect metallization materials as they meet most of the very-large-scale integra- tion (VLSI) device metallization requirements. How- ever, for the development of ultralarge-scale integra- tion (ULSI), the electrical resistivity and susceptibility to electromigration of Al and its alloys are of concern. It is becoming increasingly important to decrease signal propagation-delay times on integrated circuits (ICs). One major component of the delay time is resistivity, which depends on the choice of metallization. The resistivity of Al alloys (about 3 kfi cm) is higher than desired for fast ICs. Also as interconnect linewidths become narrower and the current densities that are carried through these conducting lines become correspondingly higher, Al metallurgy becomes in- adequate owing to thermal stress voiding and electromigration.

For the reasons presented above, a new metallization material with low resistivity and higher resistance to electromigration is needed to replace aluminum. The high electromigration resistance and low resistivity of copper (1.7 ~fl cm) make it an attractive candidate [l]. However, copper suffers from a few

0040.60901951$09.50 0 1995 Elsevier Science S.A. All rights reserved SSDI oO4(r-6(190(94)OSH05-9

drawbacks for application in metallization schemes. Copper exhibits poor adhesion when deposited on silicon dioxide (SiO,) [2]. This is due to the inability of copper to reduce SiO,, which causes poor bonding at the interface. Copper also has a high diffusivity into Si and into SiO, under bias thermal stress [3]. This high diffusivity causes problems if copper is placed in direct contact with silicon; copper diffuses into and alters the characteristics of the silicon. Unlike aluminum, copper lacks the property of self-passivation. This deficiency makes copper susceptible to oxidation during processing [4]. Cu also reacts with most noble and near-noble metals, such as Au, Ni and Al at temperatures as low as 300 “C [S].

To overcome the thermal stability problems men- tioned above, associated with Cu-based metallization, it has been proposed to employ a Cu-refractory metal (such as Ti, Cr or Nb) alloy system [6]. The refractory metal chosen must reduce VLSI-ULSI interlevel dielectrics, such as SiO, and also form a stable surface nitride. The reaction of the refractory metal with SiO, vastly improves the adhesion of copper to the dielec- tric [7], while the nitridation reaction encapsulates the copper with a nitride layer to protect it from oxidation during subsequent ULSI processing [8]. A major concern, however, with the Cu-refractory alloy sys-

200 D. Adams et al. I Thin Solid Films 262 (1995) 199-208

terns is the formation of copper-refractory intermetal- lies which may affect the resistivity of the copper interconnect material. In the range of most processing temperatures, refractory metals such as Cr, Ta and V do not react with copper. However, Ti and Zr react with Cu at 325 “C and 400 “C respectively [9].

Among diffusion barriers, titanium nitride is of increasing interest owing to its low electrical resistivity, refractory nature (which implies a high melting point), chemical inertness and extreme hardness [lo]. It has been shown that TIN is effective as a diffusion barrier in Cu/TiN/Si and Cu/TiN/Al structures up to 500 “C [ll, 121. There are various methods for forming TiN films for diffusion barrier applications; such methods include nitridation of pure Ti films in a conventional resistance-heating furnace or in a rapid thermal processor using lamp heating in N, or in an ammonia (NH,) ambient. Other preparation methods include reactive sputtering of Ti in an N, ambient and chemical vapor deposition (CVD). TiN barrier formation by reactive sputtering is a simple process. For this reason, reactively sputtered TIN is widely used as a diffusion barrier for fabrication of metal-oxide-semiconductor IC devices. The technique, however, is very process intensive from a device fabrication point of view (number of process steps, cost etc.) and is inclined to increased defect densities [13]. CVD pro- cesses utilize high processing temperatures; therefore CVD films possess greater film stresses and correspondingly are more susceptible to microcrack failure [14]. Such problems make TiN formation by rapid thermal or conventional resistance furnace heating of titanium films in N, or NH, a more desirable process. However, researchers have found limits to the quality of the resulting titanium nitride and its thickness. The degree of nitridation of the titanium can vary ap- preciably as a function of its depth, resulting in etch problems [15]. It also has been shown that TIN films formed in a single-step rapid thermal nitridation process in an NH, ambient (regardless of its thickness and the nitridation process) will not make a metal- lurgically reliable diffusion barrier [ 131.

In the present investigation we isochronally annealed Cu-Ti alloys in a conventional resistance- heated furnace in a flowing NH, ambient to form TiN-encapsulated Cu layers. The TIN passivation layers showed significant thermal stability and chemical inertness with respect to Cu. We also investigated the encapsulation of Cu layers resulting from annealing of Cu-Cr alloys in an NH, ambient. The choice of chromium stems from the fact that it does not form any intermetallic compounds with copper and has a limited solid solubility, 0.89 at.% Cr in Cu (f.c.c.) and Cu has a negligible solid solubility in Cr (b.c.c.) under equilibrium conditions [16, 171. It has been shown that aging copper alloys containing small amounts of Cr at

477 “C causes fine Cr particles to be dispersed homogeneously throughout the copper matrix, increasing the hardness [18]. Other possible improve- ments besides mechanical stability are electromigration lifetime, resistance to corrosion and adhesion to SiO, and other dielectrics [19]. The integrity of the Cr nitride formed from these NH, annealed Cu-Cr alloys was evaluated for use as diffusion barriers for Cu and Al.

2. Experimental details

Alloy films consisting of 163 nm Cu-(27at. % Ti) and 192 nm Cu-(26at.% Cr) were codeposited by electron beam evaporation onto thermally grown SiO, (110 nm) on Si( 100) substrates. The base and operat- ing pressures of the deposition system were about 5 x 1O-8 Torr and 5 x lo-’ Torr respectively. The stoichiometry and thicknesses of the as-deposited samples were determined by Rutherford backscattering spectrometry (RBS). Samples were annealed for 30 min at temperatures ranging from 400 to 700 “C in a Lindberg single-zone quartz-tube furnace in a flowing electronic grade (99.99 + %, with less than 33 mol. ppm H,O and less than 10 mol. ppm 0, + Ar) NH, ambient to form the refractory metal nitrides. A 45 nm Al film was deposited after the nitridation step, whereupon a second anneal was performed in flowing N,-5%H, forming gas. The Al depositions were performed in a bell-jar thermal evaporator with a pressure of 10e6 Torr during deposition. Before each anneal the annealing chamber was evacuated to about 10 mTorr followed by a 2.5 min purge with NH, or forming gas. This sequence was repeated twice with a final 20 min purge. The annealed samples were also allowed to cool in flowing gas before removal from the chamber. Flow rates of about 8 L min-’ were main- tained during the anneals.

Free surface as well as interfacial reactions were analyzed by RBS and Auger electron spectroscopy (AES). Microstructural analyses of the samples were performed by transmission electron microscopy (TEM). Sheet resistance values of the samples were determined from four-point-probe measurements. The RBS analysis was performed using a 1.7 MV Tandem accelerator. The backscattering angle was 170 “C and the total accumulated charge was 20 ).LC. We utilized the computer simulation program RUMP for simulation and interpretation of RBS spectra [20]. A 2 MeV He*+ beam energy was used for all RBS analyses and the samples were tilted at 7”. All anneals in this study were performed for 30 min in flowing NH, or forming gas. Unless otherwise stated, the conditions men- tioned above will be implied. The Auger instrument used was a Perkin-Elmer PHI stationary electron

D. Adams et al. I Thin Solid Films 262 (1995) 199-208 201

beam thin film analyzer which was controlled by a desktop microcomputer using a PC-137 interface. A 5 keV primary electron beam of 3 PA was used for excitation of the Auger spectra. The diameter of the beam was 2.50 km. Depth profiles were acquired by sputtering with 2 keV Ar ions, directed at 30” to the sample surface normal and rastered over a few mil- limeters. It should be noted that no attempt was made to separate the overlapping Ti(LMM) and N(KLL) transitions near 380 eV. The Si signal was monitored in the energy window 76-102 eV. Cross-sectional samples for TEM analysis were prepared by using mechanical dimpling followed by argon ion beam milling.

3. Results

3.1. Free surface and interfacial reactions

3.1.1. Cu-Ti system. Figure 1 compares the Ruther- ford backscattering spectra obtained from the as-deposited alloy with those from alloys annealed at two different temperatures. Since the atomic masses of Cu and Ti are close to each other, their backscattering spectra overlap partially. Simulation of the RBS data of the as-deposited sample indicates a composition of Cu-(27at.% Ti) and an alloy thickness of 163 nm. The sample annealed at 450 “C differs from the as-deposited sample only by the presence of a thin surface oxide layer of about 10 nm. The increased surface Ti peak suggests that, at this temperature, Ti segregates to the free surface, presumably to react with oxygen and nitrogen from the ambient. This out-diffusion of

Energy (MeV)

04 0.6 0.8 1.0 1.2 1.4 1.6 I I I I I I I I 1

80 -As-depositea - -45092, 30 ml”

Surface Ti

- -5OOT. 30 min TiO. TN.(O) j

-0 .; 60

m

: .- 0 40

E ;;

Z 20

Surface Oxide

l&O 260

Channel

Fig. 1. 2 MeV He” spectra obtained from as-deposited Cu-

(27at.% Ti) alloy on SO, and from the alloy annealed at 450 and

500 “C for 30 min in NH,. Since the atomic masses of Cu and Ti are

close to each other, their backscattering spectra overlap.

Ti makes the subsurface layer of the alloy Cu rich. No segregation or interfacial reaction of Ti at the alloy- SiO, interface is observed. However, the sample annealed at 500 “C clearly shows the appearance of two separate Ti signals, resulting from the segregation of Ti to the free surface and also to the alloy-SiO, interface. The titanium at the free surface is labeled Surface Ti and that at the alloy-SiO, interface Interfa- cial Ti (Fig. 1). Owing to the segregation of the Ti to the free surface upon annealing, the surface position of Cu shifts to lower energies. At 550 “C and higher temperatures the integrated surface-Ti signal decreases compared with the interfacial Ti signal; this can be seen from the Rutherford backscattering spectra shown in Fig. 2. This increase in interfacial Ti is possibly due to reaction between Ti and the underlying SiO, which competes favorably with the surface reaction in this temperature regime. At 650 “C, less than 10% of the Ti segregates to the free surface as indicated by the small surface Ti peak in Fig. 2. Most of the Ti segregates to the alloy-SiO, interface to react with the Si and 0 freed upon the dissociation of the SiO,. Simulation of the RBS data shows that the Cu-rich layer contains less than 4% of residual Ti.

In Fig. 3, AES depth profiles obtained from Cu-Ti samples nitrided at 550 and 600 “C are shown. Al- though the N(KLL) and Ti(LMM) signals near 380 eV overlap, presence of Ti nitride shows up as an en- hancement of the Ti(380 eV) peak. The profile based on this 380 eV peak is indicated by the label Ti + N in Fig. 3. The oxygen and Ti + N profiles in the time interval O-20 min (Fig. 3(a)) suggest the existence of a two-layer surface structure. This two-layer surface

Energy (MeV)

0.4 0 6 0.8 1.0 12 1.4 1.6

Channel

Fig. 2. 2 MeV He” spectra obtained from Cu-(27at.% Ti) alloy

nitrided at temperatures of 550, 600 and 650 “C for 30 min in flowing

NH,. The smaller surface Ti signal at 650°C is related to the stronger interaction between Ti and SO, substrate at this tempera-

ture.


100 80 60 40 20 0

Sputter time (min)

$ 0 5

Cu(Ti) =!

2

Si SiO 2 P “, U * 9

b) 6OO’C. 30 min, NH,

_c

.L?

2

a

d

100 80 60 40 20 0

Sputter time (min)

Fig. 3. Auger depth profiles of Cu-(27at.% Ti) isochronally an-

nealed for 30 min at (a) 550 “C and (b) 600 “C in NH,. The surface

Ti nitride is labeled Ti + N due to the overlap of the N(KLL) and

Ti(LMM) signals near 380 eV.

structure consists of a Ti-oxide layer (labeled TiO,) present below a Ti nitride layer (labeled TiN,(O)). In addition, a thin native oxide layer (not shown in the schematic diagram) is apparent from a second rise in the oxygen signal at the surface of the structure. Note that at this temperature the Ti signal is reduced substantially in the Cu(Ti) layer, suggesting a minimal (4% or less) presence of residual Ti in the Cu. The 550 “C anneal Auger electron spectra also clearly show the presence of an additional oxygen peak, immediately below the Cu-rich layer in Fig. 3(a). The source of the oxygen is the SiO, substrate which is being reduced by the Ti. As the oxide is dissociated, silicon reacts with Ti and forms a Ti silicide (labeled Ti,Si, in the schematic diagram) adjacent to SiO, and a Ti oxide (labeled TiO,) on top of the silicide. The presence of this silicide layer manifests itself by the Si peak located near 67 min. Figure 3(b) shows AES depth profiles obtained from the sample annealed at 600 “C. Com-

pared with the 550 “C anneal, the surface nitride (TiN,(O)) layer seems to be more uniform in composition and contains less oxygen. However, at this temperature, significant reaction between Ti and the underlying SiO, substrate is observed, as indicated by the larger oxygen and silicon signals located at 60 min and 67 min respectively.

Figure 4(a) shows a cross-section transmission electron micrograph of the Cu-Ti alloy after annealing at 500 “C. Various layers are indicated on the micrograph. An oxygen-rich surface reaction layer of about 10 nm is labeled TiO,/TiN,(O). Below this surface layer is a Cu layer with some residual Ti, indicated as Cu(Ti). The thickness of this layer is about 170 nm. A Ti,Si,/TiO, structure of total thickness about 3-4 nm is located adjacent to the SiO, substrate. These layers are more clearly discernible in the case of 600 “C anneals. The multilayer structure obtained after a 600 “C anneal (Fig. 4(b)) is similar to that of the 500 “C anneal, but with well-developed interfacial reaction layers. The total thickness of the surface reaction layer is about 20 nm and consists of TiN,(O) and TiO, phases. The grain sizes in these layers of the samples annealed at 500 “C and 600 “C are about 20 nm and about 30 nm respectively. The reaction layers at the Cu-SiO, interface were identified as Ti,Si, and TiO, with a total thickness of about 20-25 nm. Electron diffraction and high resolution TEM were utilized for phase analysis. The interface between these layers and the SiO, is rough, on a scale of about 5 nm.

3.1.2. Cu-Cr system. Nitridation results for the Cu- (26at.% Cr) alloy annealed at various temperatures are compared with results from the as-deposited sample in Fig. 5. No difference is observed between the as-deposited and 450 “C annealed sample. How- ever, after a 500 “C anneal a significant amount of Cr has segregated to the surface. The Rutherford backscattering spectra indicate that the Cr segregates only to the free surface to form a Cr-nitride layer. Within the detection limit of RBS, no interfacial reaction was observed at the alloy-SiO, interface. At 600-650 “C it seems that most of the Cr has segregated to the free surface (Fig. 6). Simulations of the RBS data of the samples nitrided at these temperatures indicate that the Cu-rich layer contains less than 2% residual Cr. By comparing the results of the two alloy systems, it is evident that thicker nitride layers are formed in the Cu-Cr system than in the Cu-Ti system and also that the oxygen composition in the surface layers is lower in Cu-Cr than in Cu-Ti.

The AES depth profile for the Cu-(26at. % Cr) sample annealed at 550 “C (Fig. 7(a)) shows the formation of a relatively thick surface nitride layer (CrN,(O)) with much less oxygen than in the nitrided Cu-Ti sample. An interesting feature of the oxygen


25 nm PI

Fig. 4. Cross-section TEM bright-field images of the multilayer

structure of the Cu-(27at.% Ti) alloy annealed for 30 min at (a) 500

“C and (h) 600 “C in NH,. A TiO,$ /Ti,Si, bilayer is located at the

Cu-SiO, interface.

profile in this sample is the presence of a peak at the Cr-nitride-Cu interface, labeled as Cr(0) in the schematic diagram. The Cr-nitride layer resulting from the 600 “C anneal (Fig. 7(b)) seems to be more uniform than for the 550 “C anneal. It is also evident that at this temperature the original Cu-(26at.% Cr) sample has been almost completely dealloyed.

Energy (IJeV!

0.4 0.6 0.8 1 0 ’ 2 ( .4 1.6

I I

80

t

-As-depostted - -45O'C. 30 min - -5OOT 30 min

5C 103 ’ 50 2CS 253 300 I- L, q n -, 5

Fig. 5. Rutherford backscattering spectra obtained from as-depos-

ited Cu-(26at.% Cr) alloy and from the alloy annealed for 30 min at

temperatures of 4.50 and 500 “C. A 2 McV He” beam energy was

USd.

0.4 0.6

Energy (VI&‘)

0.8 1 0 1.2 i4 ’ .6

80

3

Fig. 6. Rutherford backscattering spectra of the Cu-(26at.% Cr)

alloy after nitridation at 550, 600 and 650 “C for 30 min in NH,

ambient. A 2 MeV He” beam energy was used. The 600 and 650 “C

spectra are almost identical, showing nearly complete dealloying of

Cr.

Furthermore, at this temperature no reaction at the Cu(Cr)-SiO, interface is observed.

Figure S(a) presents a transmission electron micrograph obtained from the Cu-(26at.% Cr) alloy nitrided at 500 “C for 30 min in flowing NH,. Various layers are indicated on the micrograph. A Cu(Cr) layer of thickness about 160 nm is present above the SiO,. Grain sizes in this layer vary from about 10 to 30 nm, with the grains tending to be equiaxed. On top of the Cu(Cr) layer is an additional layer labeled


v, SiO ? 2 Cu(Cr) l-

9 CW(O)

I

100 80 60 40 20

Sputter Time (min) (4

)i

I CU

/Si rO I

\- ---,

100 80 60 40 20 0

Sputter Time (min)

Fig. 7. Auger depth profiles of Cu-(26at.% Cr) alloy nitrided for 30 min at (a) 550 “C and (b) 600 “C in NH,. The peak-to-peak data shows no sign of interfacial reaction between Cr and SiO, at these temperatures.

CrN,(O). This is the Cr-nitride layer formed by the nitridation process. The layer is about 33 nm thick. Cross-correlation with results from Auger spectroscopy indicates the presence of an oxygen-rich region between the CrN,(O) and Cu(Cr) layers. This oxygen- rich layer is indicated in Fig. 8(a) by the label O-R. Figure 8(b) presents a cross-section transmission electron micrograph obtained from Cu-(26at.% Cr) alloy nitrided at 600 “C for 30 min in NH,. Once again, various layers are indicated on the micrograph by the labels CrN,(O), O-R, Cu(Cr) and SiO,. Identification of the layers is the same as discussed in the context of Fig. 8(a). Grain sizes in the Cu(Cr) layer vary from about 20 to 45 nm, with the grains tending to be equiaxed. The Cu(Cr) layer is 135 nm thick. On top of the Cu(Cr) layer, Cr nitride (about 36 nm thick) is present, resulting from the nitridation process. Once again, cross-correlation with results from AES indicates the presence of an oxygen-rich region between

lb)

Fig. 8. Bright-field transmission electron micrographs showing the multilayer structure of the Cu-(26at.% Cr) alloy annealed for 30 min at (a) 500 “C and (b) 600 “C in NH,. The different layers are separated by well-defined interfaces.

the CrN,(O) and the Cu(Cr) layers (labeled O-R in Fig. 8(b)). At both temperatures the different layers are separated by well-defined smooth interfaces. In contrast with the Cu-Ti alloy (Fig. 4) no interfacial layer is visible between Cu and SiO, in both samples after annealing at these temperatures. These results are in agreement with the RBS data presented earlier, which suggested no interfacial segregation or reaction of the Cr with SiO,.

3.2. Electrical characterization

Sheet resistances of the samples nitrided for 30 min in an NH, ambient at different temperatures were measured. The nitrided sample of the Cu-Ti alloy annealed at temperatures higher than 4.50 “C consists of a dual Ti-oxide/Ti-nitride surface layer, a Cu-rich subsurface layer and a Ti-silicide/Ti-oxide bilayer at the alloy-SiO, interface. The resulting structure of the

D. Adams et al. I Thin Solid Films 262 (199-Y) 199-20X 205

TemDerature i”C) 0

Fig 9. Sheet resistance as a function of nitridation temperature for

both Cu-(27at.% Ti) and Cu-(26at.% Cr). The as-deposited

sample temperature was taken as 25 “C.

Cu-Cr alloy after nitridation is a Cu/Cr-nitride bilayer structure on the SiO, substrate. Fig. 9 shows the sheet resistance as a function of nitridation temperature for both the Cu-Ti and the Cu-Cr alloy systems. The as-deposited Cu-Ti and Cu-Cr samples have values of about 12.69 R/O and 7.7 R/Cl respectively. For purposes of this plot, the temperature of the as- deposited samples was assumed to be 25 “C. The resistance of the Cu-Ti alloy decreases to a minimum of about 1.26 SL/O, after nitridation of the samples in the temperature range 450-550 “C for 30 min. It then remains almost constant in the interval 550-650 “C. At temperatures higher than 650 “C, the resistance starts to increase again. The Cu-Cr alloy system in general exhibits lower sheet resistance values over the temperature range considered. As in the Cu-Ti system, this system also reaches its minimum sheet resistance at about 550 “C. This minimum value (about 0.4 Q/O) remains constant up to anneal temperatures as high as 700 “C. This is in contrast with the Cu-Ti system which shows a drastic increase in resistance at temperatures higher than 650 “C.

3.3. Evaluation of diffusion barriers

Rutherford backscattering spectra of the Cu-Ti nitrided samples (550 “C for 30 min in an NH, ambient) after Al deposition (as prepared) and subsequent annealing at 400 and 500 “C for 30 min in N,-5%H, are presented in Fig. 10. The results suggest no difference between the as-prepared sample and that annealed at 400 “C for 30 min. This implies that the barrier is still intact at this temperature. However, simulation of the data obtained for the 500 “C. 30 min anneal confirm the presence of Cu at the

Energy (MeV)

1 2 : 4 16

I

t

-As-prepared - -400 c, 30 VI” - -530°C. 3c rr,n

Ti a: Fnterfoce

180 200 220 240 360 Channel

Fig. 10. 2 MeV He” Rutherford backscattering spectra of the

Cu-(27at.% Ti) nitrided sample after Al deposition (as prepared)

and subsequent annealing at 400 and 500 “C for 30 min in NL-

5%H2. The SiOz and Si substrate signals are not shown. The

Ti-nitride barrier fails at 500 “C.

surface, which indicates failure of the Ti-nitride diffusion barrier.

In Fig. 11 the effectiveness of Cr nitride as a diffusion barrier for Al and Cu is demonstrated. The as-prepared structure was obtained by nitriding a Cu- (26at.s Cr) alloy at 600 “C for 30 min in a flowing NH, ambient, followed by depositing an Al capping layer 45 nm thick. No difference is observed between the as-prepared and the 500 “C, 30 min annealed

Energy (MeV)

1.1 1.2 1.3 1.4 100

1.5 1.6 1.7

I I I I I 1

80

u a, .- F 60 -0

E .- 0 E 4o b Z

20

0

-As-prepared Cr-nitride Barrier

- -500 C. 30 min - -6OO’C, 30 min I

Al

, I I I I I 180 200 220 240 260 280 300

Channel

Fig. 11. 2 MeV He” Rutherford backscattering spectra of the

Cu-(26at.% Cr) alloy nitrided sample after Al deposition (as

prepared) and subsequent annealing at 500 and 600 “C for 30 min in

N,-5%H,. The Si and SiOz substrate signals are not shown. The

Cr-nitride barrier is stable up to 600 “C.

206 D. Adams et al. I Thin Solid Films 262 (19Y5) 199-208

samples suggesting that the Cr-nitride barrier is still intact at this temperature. The RBS results indicate that the Cr-nitride layer is effective as a diffusion barrier in the Al/CrN,(O)/Cu structure for temperatures lower than 600 “C. After 600 “C annealing, Cu was detected at the surface, implying failure of the Cr nitride as a diffusion barrier.

Discussion

4.1. Free surface and interfacial reactions

The results obtained from annealing the Cu- (27at.% Ti)/SiO, /Si structure in an NH, ambient at different temperatures for 30 min indicate that Ti segregates out of the Cu to both the free surface and the alloy-SiO, interface. A two-layer Ti-nitride/Ti- oxide surface layer is formed at the free surface owing to the reaction of Ti with the NH, ambient. It is believed that the formation of the surface layer proceeds as follows. At the onset of the annealing process, oxygen and nitrogen from the ambient react with Ti which has diffused to the free surface, to form Ti-oxide/Ti-nitride phases. The large negative heat of formation of Ti oxide (222-365 kJ mol-‘) compared with that of Ti nitride (169 kJ mall’) makes formation of the former more thermodynamically favorable. However, the composition and order of the reaction layers suggest that the reactions are not driven by thermodynamic factors alone, but possibly by a combi- nation of thermodynamics and kinetics. It seems that, once the nitride layer starts to form, it impedes the in-diffusion of oxygen, as observed by the minimum in the oxygen signal in the surface region. It is speculated that this phenomenon is due to the low solubility and diffusivity of oxygen into the nitride phase [13]. The thickness and uniformity of the Ti-nitride layer increase with increasing temperature as observed in Fig. 3(b). Hoshino et al. [21] showed that annealing Cu-Ti in a nitrogen ambient at 800 “C can result in the formation of a TIN layer on the copper surface. It is evident that titanium nitridation of the Cu-Ti alloy in NH, occurs at markedly lower temperatures than in N,. The relatively low temperature required to form TIN in an NH, ambient compared with N, could be due to the difference between the thermal decomposi- tion energies of the two gas molecules. The decompo- sition energies of N, and NH, are 942 kJ mall’ and 432 kJ mall’ respectively [22].

Our results indicate that significant reaction occurs between Ti and SiO, at temperatures as low as 500 “C. The reaction starts with the dissociation of SiO,; the released oxygen and silicon react with Ti to form a bilayer structure at the alloy-SiO, interface. A Ti-rich silicide layer is formed adjacent to the oxide substrate

and a Ti-oxide layer forms on top of the silicide. Cross-section TEM analysis showed that the reaction results in the formation of a TiO,/Ti,Si,lSiOz stack at temperatures in the range 500-700 “C. Russell et al. [23] have shown that, as the silicide formed from annealing SiO,/Ti in vacuum grows, the concentration of 0 in TiO, increases with the reaction ceasing at w = 1.2. This interfacial reaction is simultaneously accompanied by the free-surface reaction to form a TiN,(O) layer. It has been previously reported that the reaction between Ti and SiO, in Ar starts at approximately 400 “C and that loss of SiO, becomes significant above 500 “C [24]. Furthermore, a strong interaction occurs between Ti and SiO, at 700 “C and above.

The nitridation results of the Cu-(26at.% Cr) alloy annealed for 30 min at different temperatures in a flowing NH, ambient indicate that Cr segregates only to the free surface, to form a Cr-nitride layer. Within the detection limits of the experimental techniques used, no reaction was detected at the alloy-SiO, interface. Chromium does not form any intermetallic compounds with copper and has essentially no solid solubility at room temperature in copper. The pre- domination of the free-surface reaction in Cu-Cr alloys is then responsible for the formation of thicker nitride layers than in the case of Cu-Ti alloys. The nitridation process occurs by outdiffusion of Cr from the Cu-Cr alloy to react with NH, in the ambient.

4.2. Electrical properties of nitrided samples

Interpretation of the sheet resistance behavior of the nitrided Cu-Ti and Cu-Cr alloys will be based on two concepts. Firstly, Matthiessen’s rule [25] applies, in that the resistivity of pure metals and dilute alloys can be separated into two components:

P c total Pphonon + Presidual (1)

where ~~~~~~ is the total resistivity, pphonon is the temperature-dependent electron-phonon contribution and presidua, is the residual resistivity. The residual contribution arises from electron scattering by im- purities and structural imperfections such as grain boundaries, dislocations and vacancies. In the Cu- based alloys, the presence of refractory-metal solute atoms in the Cu matrix will lead to additional electron scattering centers and hence a higher resistivity than in the case of pure copper. However, if the Cu can be completely dealloyed upon thermal annealing, without the formation of any intermetallic compounds, the resistivity should decrease to a value close to that of pure Cu.

Secondly, the sheet resistance of a bilayer structure in which current flows parallel to the bilayer interface can be calculated by assuming that the two layers act


as parallel resistors, of resistances R, and R, [26]. The total electrical resistance of the film is then given by

1 1 1 R,,,,=R,+cc, (2)

In the structures formed after nitridation of the Cu-Ti and Cu-Cr. it is assumed that the Cu-rich layer has the lowest resistance (R,,) and it is also the thickest layer. Therefore, from Eq. (2) l/R,, > l/Rpassivat,on

CRparsivatlon is resistance of the passivation layers) and

1 ‘KM,, = l~R,U or Rtota, = R,, (after annealing). The variation in the sheet resistance with nitridation

temperature for the Cu-Ti in Fig. 9 shows a significant decrease in the interval 450-550 “C. It is believed that this decrease is related to the dealloying of the Cu that occurs in this temperature range. The AES results (Fig. 3) indicate that at 550 “C most of the Ti has segregated to the free surface and interface, leaving behind an almost pure Cu layer. At temperatures of 550 “C and higher the interfacial reaction with SiO, becomes dominant (see Fig. 2(b)). It is speculated that this process results in an effective increase in the thickness of the conductive layer due to the formation of the Ti,Si, layer, which manifests itself by the lower sheet resistance values. However, at temperatures higher than 650 “C, the sudden increase in sheet resistance may be due to diffusion of Cu into the Ti,Si,/TiO, layers and the subsequent formation of intermetallic compounds.

Both AES and RBS results show that, in the Cu-Cr alloy, Cr segregates only to the free surface during thermal annealing to form thick surface nitride layers. The variation in sheet resistance with nitridation temperature seems to reach a minimum value near 550 “C. It further shows that the resistance does not change with increasing temperature after the minimum value is reached. This can be explained by the existence of an almost completely dealloyed Cu subsurface layer at temperatures as low as 550 “C. Shih et al. [26] have shown that in Cu-based bilayer systems (involv- ing metals that are insoluble in Cu, such as Cr and Co) the resistivity of the Cu film remains unchanged or even decreases slightly after annealing.

4.3. Evaluation of diffusion barriers

In this section we discuss the effectiveness of TiN,(O) and CrN,(O) barriers formed by annealing Cu-Ti and Cu-Cr alloys in an NH, ambient at various temperatures for 30 min. The term “as prepared” will refer to the nitrided alloy with Al deposited on it. Therefore the integrity of Ti nitride as a diffusion barrier between Al and Cu was evaluated in this study. Because the nitrided samples were exposed to air for several hours prior to Al deposition, the Ti- nitride grain boundaries can be considered to be

stuffed with oxygen. It has been observed in the past that the TIN films form a better barrier to the intermixing of aluminum and silicon when exposed to air prior to Al deposition [27-291 or when the TiN contains significant amounts of oxygen [30]. Li et al. [31] have also shown that TiN with oxygen incorpo- rated in the lilm is more effective in blocking Cu inward diffusion in both Cu/TiN/Al and CulTiN/ silicide systems.

RBS analyses indicate no difference between Cu/ TiN,(O)/Al annealed at 400 “C for 30 min in N,- 5%H, and as-prepared sample. However, the results obtained for the 500 “C anneal showed the presence of Cu at the surface, which suggests failure of the Ti- nitride diffusion barrier. The results further suggest that the Ti-nitride film dissolves as a result of the heat treatment. Since Al can be nitrided and Cu can react with Ti to form compounds when these elements come into contact with Ti nitride during heat treatment, it is believed that the formation of Al,Ti, Cu,Ti and AIN compounds were primarily responsible for the dissociation of the TIN barriers.

In the case of Cu-Cr alloys, the “as-prepared” structure was obtained by nitriding the Cu-(26at.% Cr) alloy at 600 “C for 30 min in a flowing NH, ambient, followed by depositing an Al capping layer 45 nm thick. The RBS data indicated that the Cr- nitride layer is effective as a diffusion barrier in Cu/ Cr-nitride/Al structures up to 600 “C. At this temperature, Cu was detected at the surface. The CrN,(O) diffusion barrier does not dissolve as an effect of the heat treatment, although it allows migration of Cu. It is speculated that the diffusion occurs along grain boundaries and possibly microvoids in the barrier layer.

5. Conclusions

The results of this investigation show that an encapsulated Cu layer can be formed by annealing either a Cu-(27at.% Ti) or a Cu-(26at.% Cr) alloy at 550-700 “C for 30 min in flowing NH,. A TiO,/ TIN,(O) passivation layer structure is formed over the Cu-Ti alloy. Conventional depth-profiling techniques have shown that this surface structure is effective as a diffusion barrier for Al and Cu up to 500 “C. Anneal- ing of the Cu-Ti system also resulted in the formation of a Ti,Si,/TiO, bilayer at the alloy-SiO, interface. At higher temperatures it seems that this interfacial reaction is more favored than the free surface process. CrN,(O) passivation layers formed by nitridation of the Cu-Cr alloys were stable up to 600 “C against Cu and Al interdiffusion. No interfacial reactions or segregation of Cr at the alloy-SiO, interface were detected. Evaluation of the electrical properties

208 D. Adams et al. 1 Thin Solid Films 262 (1995) 199-208

showed that the Cu-Cr nitrided samples have lower sheet resistances than the Cu-Ti nitrided alloys. Both systems show a decrease in sheet resistance with increasing nitridation temperature. However, for the Cu-Ti system the resistance increases again at temperatures higher than 650 “C; the increase may be due to intermetallic compound formation.

Acknowledgments

The authors would like to thank Harland Tompkins (Motorola, Mesa, AZ) for the AES analysis and Moon J. Kim (center for high resolution electron microscopy, Arizona State University) for assistance with the cross-section TEM analysis. We also acknowl- edge the financial support provided by the National Science Foundation (L. Hess, DMR-9307662, and L. Salmon, ECS-9410399).

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Passivation of Cu via refractory metal nitridation in an ammonia ambient