J O U R N A L O F M AT E R I A L S S C I E N C E : M AT E R I A L S I N E L E C T RO N I C S 1 0 ( 1 9 9 9 ) 2 8 5 ± 2 9 0

Metalorganic chemical vapor deposition ofaluminum from tetramethylethylenediamine alane

DAE-HWAN KIM, MAN-YOUNG PARK, SHI-WOO RHEELaboratory for Advanced Materials Processing (LAMP), Department of ChemicalEngineering, Pohang University of Science and Technology (POSTECH), Pohang,790-784, KoreaE-mail: srhee@postech.ac.kr

Tetramethylethylenediamine alane (TMEDAA) was synthesized by ligand displacementreaction of dimethylethylamine alane (DMEAA) with N,N,N'N'-tetramethylethylendiamine(TMEDA), and the chemical vapor deposition of aluminum ®lm from TMEDAA in thetemperature range of 140±260 �C has been studied. The maximum deposition rate of Al ®lmfrom TMEDAAwas 140 nm/min at 210 �C and the apparent activation energy over a substratetemperature range of 140±210 �C is about 58.6 kJ/mol. Al ®lms were deposited on TiN/Sisubstrate and electrical resistivity values in the range 5±35 mOcm were obtained. Theincorporation of carbon and oxygen, and surface roughness were increased as the substratetemperature was increased. The Al ®lms with a preferred orientation of (1 1 1) were obtainedover a wide range of substrate temperature.

1. IntroductionAluminum has been used widely as a conducting

material in the fabrication of integrated circuits [1]. So

far, most commercial Al ®lms have been deposited by

physical vapor deposition (PVD), i.e. sputtering or

evaporation. However, chemical vapor deposition

(CVD) usually gives more conformal coverage on the

substrate surface and allows a reactor design for multi-

wafer processing to give high throughput. For this reason,

CVD of Al has drawn a great deal of attention as a

promising process for next-generation metallization.

For the deposition of Al ®lms, alkyl aluminum

precursors such as triisobutylaluminum (TIBA) [2],

trimethylaluminum (TMA) [3], and dimethylaluminum

hydride (DMAH) [4, 5], and amine-alane adducts such as

trimethylamine alane (TMAA) [6], triethylamine alane

(TEAA) [7], and dimethylethylamine alane (DMEAA)

[8±11] have been studied. Particularly, amine-alane

adducts, which have no direct Al±C bond, have been

found to yield high purity Al. For example, DMEAA, an

adduct of alane and dimethylethylamine (DMEA), is the

most recently introduced member of the amine family of

precursors. Its relatively high vapor pressure (200 Pa) at

room temperature and its ability to deposit carbon

contamination-free ®lms, combined with the advantages

of being a liquid have made it the precursor of attention

recently.

However, DMEAA slowly degrades during storage

and forms aluminum precipitation and is dissociated into

dimethylethylamine and alane even at room temperature

in the gas phase [12]. Changes in the electron donating

ability of the amine, the Lewis-base, are likely to

enhance the stability of this amine-alane adduct and we

adopted a new amine-alane adduct precursor, tetra-

methylethylenediamine alane, which has a greater

electron donating ability of the amine. In this research,

the deposition and properties of aluminum thin ®lm with

tetramethylethylenediamine alane (TMEDAA) have

been investigated.

2. Experimental proceduresTMEDAA preparation has been described in detail

elsewhere [13±15] and in this work, the synthesis of

TMEDAA was achieved in one step through the

following reaction.

H3AINMe2Et�Me2N(CH2�2NMe2

?H3AIMe2N(CH2�2NMe2 � NMe2Et

A cold wall, low-pressure chemical vapor deposition

(LPCVD) reactor was fabricated to deposit Al ®lm from

TMEDAA on a small scale. The schematic diagram of

the Al MOCVD (metalorganic chemical vapor deposi-

tion) reactor is shown in Fig. 1. TMEDAA was loaded

into the precursor bottle in N2 atmosphere and was

carried into the reactor by H2 carrier gas at a ¯ow rate of

15 sccm. To supply a suf®cient amount of TMEDAA, the

precursor bottle was maintained at 90 �C and the feeding

line was maintained at 100 �C to prevent precursor

condensation. The base pressure of this pumping system

was about 1 Pa. The susceptor was resistively heated and

the total pressure in the reactor was adjusted to 133 Pa by

the throttle valve between the pump and the reaction

chamber. The Si wafer coated with sputtered TiN was

used as a substrate. The TiN substrate was rinsed with

deionized water and then dried with N2 gas. The

deposition rate was calculated from the increase in the

0957±4522 # 1999 Kluwer Academic Publishers 285

weight of the substrate and con®rmed by cross-sectional

scanning electron microscopy. The resistivity was

determined by the four-point probe method (Chang-

Min SR1000). An X-ray diffractometer (MAC Science

Co. M18XHF) with CuKa radiation was used to

determine the crystal structure and preferred orientation

of Al thin ®lm. The ®lm composition was evaluated by

Auger electron spectroscopy (AES, Perkin-Elmer PHI

600) and the surface morphology of Al ®lm was observed

using a scanning electron microscope (SEM) and an

atomic force microscope (AFM, Park Scienti®c

Instruments Autoprobe-CP).

3. Results and discussionFrom the 1H-NMR (nuclear magnetic resonance) and

mass spectrum measurement, it was con®rmed that the

compound was the 1:1 amine-alane adduct of TMEDA

and alane. Davison and Watrik [13] have reported that

TMEDAA is a stable, non-pyrophoric, white solid with a

high vapor pressure (200 Pa at 99 �C, 1400 Pa at 119 �C)

and, unlike other AlH3 adducts, showed no tendency to

decompose over a 24-h period at 133 �C. It is believed

that TMEDA has a greater af®nity of the N-donor to the

alane with a tendency to form ®ve-co-ordinate alane-

amine adducts [16].

Fig. 2 shows plausible molecular structures of

TMEDAA. The molecular structure and polymeric

nature of TMEDAA in the solid phase has been reported

earlier [13, 17]. Gaseous TMEDAA is known to be

dimerized. Davison and Watrik had proposed a

hydrogen-bridged dimeric structure in the gas phase in

which TMEDAA was thought to adopt a bidentate hexa-

co-ordination mode (Fig. 2c) [13], while Young and

Ehrlich [15] suggested the penta-co-ordinated dimeric

structure (Fig. 2b). It is well known that the Al±H

frequency correlates with the co-ordination number of

the aluminum atom and Al±H stretching band is known

to appear at 1709±1710 cmÿ 1 in the case of a penta-co-

ordinated structure such as tetramethylpropanediamine

alane (TMPDAA) and bis(trimethylamine) alane

(BTMAA). From the IR spectrum of TMEDAA using a

gas cell, the peak of the Al±H stretching band appeared at

1709 cmÿ 1 at cell temperature of 200 �C and it is more

likely that molecular structure of TMEDAA is cyclic

dimer (Fig. 2b) in the gas phase.

The deposition rate was measured over a substrate

temperature range of 140±250 �C at the constant reactor

pressure of 133 Pa. Fig. 3 shows the Arrhenius plot of the

deposition rate of Al ®lm from TMEDAA and that of

DMEAA is also shown for comparison [10]. As shown in

Fig. 3 the logarithm of deposition rate decreased linearly

with respect to reciprocal temperature in the range of

140±210 �C and the maximum deposition rate from

TMEDAA was about 140 nm/min at 210 �C. The

deposition rate was increased to a maximum below

210 �C and then decreased with substrate temperatures

above 210 �C. The apparent activation energy of the

deposition calculated from the slope below 210 �C was

found to be about 58:6+7 kJ/mol which is greater than

that of DMEAA (41.9 kJ/mol) below 150 �C. We believe

that this large activation energy suggests surface-

reaction-limited growth rather than mass-transfer-limited

growth at lower temperature region. Higher deposition

temperature and activation energy with TMEDAA is due

to the enhanced thermal stability of TMEDAA from the

stronger Al±N bond in the chelating structure.

At higher deposition temperature (above 210 �C), the

growth rates of TMEDAA decreased rapidly and did not

show an adjacent plateau region (mass-transfer-limited

region) which occurred in the Arrhenius plots of the

growth rate of DMAH [5]. The decrease of growth rate at

high temperature is in line with the growth result of

Figure 2 The chemical structures of TMEDAA; (a) monomeric, (b)

dimeric ( penta-co-ordinated), (c) dimeric (hexa-co-ordinated), and (d)

polymeric forms.

Figure 1 Schematic drawing of Al MOCVD apparatus using

TMEDAA. Total reactor pressure was maintained at 133 Pa and H2

carrier gas at a ¯ow rate of 15 sccm was introduced during the

deposition.

Figure 3 Arrhenius plot of the deposition rate of Al ®lm from

TMEDAA (�) and DMEAA (�) [10].

286

DMEAA [10] and is not unusual in organometallic

deposition. Recently, the fall off in the growth rate of

DMEAA at higher temperature was well explained by

the gas phase dissociation into alane and amine [12]. The

gas phase reaction of DMEAA has been investigated by

in situ Fourier transform infrared (FT-IR) spectroscopy

and it was found that DMEAA was dissociated in the gas

phase very rapidly at temperatures above 150 �C, which

resulted in the decrease of the deposition rate. TMEDAA

is also one of the family of alane-amine adducts and

shows a similar deposition tendency, and the decrease of

the deposition rate at high temperatures was probably

caused by the decomposition of TMEDAA into diamine

and alane in the gas phase. The amine group is a neutral

ligand and has a relatively weak bond with the aluminum

atom. The dissociation of gas phase alane-amine adduct

may account for the growth rate variation with this

compound. The alane (AlH3) is known to be very

unstable and reactive and is more likely to react with any

residual impurity in the gas phase to form particles or

reactive species. The particle produced in the gas phase

can be incorporated into the ®lm, leading to poor

morphology and the higher level of impurities in the

®lm with increased resistivity at higher temperature [18].

Al ®lm prepared from TMEDAA at higher temperature

showed a higher impurity level and resistivity. However,

a more detailed study on the effect of the gas phase

reaction would be needed to clarify the deposition

kinetics of TMEDAA.

The electrical resistivity of the deposited Al ®lms with

thickness of 900 nm was measured by the four-point

probe method. Fig. 4 shows the change of resistivity as a

function of the substrate temperature. The lowest value

of the electrical resistivity of Al ®lm was 4.9 mO cm,

which is larger than 2.7 mO cm of the bulk electrical

resistivity. Generally, grain morphology, ®lm thickness,

impurity incorporation, voids or porosity, and surface

roughness affect the resistivity of the thin ®lm. The

primary reason for the large resistivity change from the

bulk value at low growth temperatures is probably

electron scattering at grain boundaries. At low growth

temperatures, the ®lms showed a smaller grain size and

had lower impurity than the ®lms deposited at higher

temperatures. Films with small grains have larger grain

boundary area, which is expected to be effective as an

electron scatterer. As the substrate temperature was

increased, the resistivity gradually increased, probably

due to the increased impurity content.

Fig. 5 shows an AES spectra, taken after 100 nm

argon-ion sputtering, of Al thin ®lms with a thickness of

900 nm deposited at different substrate temperatures. The

signi®cant increase in the electrical resistivity and in the

level of impurities of the ®lm was observed at high

deposition temperatures. Incorporation of oxygen and

carbon was less than 7 at % at low substrate temperatures;

below 210 �C. Al ®lm deposited at 260 �C has oxygen

and carbon impurities at about 19 at % and 20 at %,

respectively, and it was easily stripped off due to the very

poor adhesion to the substrate, and showed a dark gray

instead of a milky color. It is reasonable to believe that at

high temperatures, the resistivity increases dramatically

due to the carbon and oxygen incorporation which leads

to the void formation and poor coalescence of the grains.

The high oxygen concentration may be ascribed to

internal oxidation after deposition of the porous ®lms or

incorporation from the residual impurity gases such as

H2O, CO2, or O2 in the CVD reactor during growth.

Carbon contamination may arise from the decomposition

of TMEDA ligand at the surface or in the gas phase.

X-ray diffraction (XRD) studies of Al ®lms con®rmed

the formation of polycrystalline aluminum. Fig. 6 shows

Figure 4 Electrical resistivity of CVD-grown aluminum ®lms with

thickness of 900 nm as a function of the growth temperature.

Figure 5 Auger electron spectrum (AES) of thin ®lms with thickness of

900 nm deposited at (a) 180 �C, (b) 240 �C, and (c) 260 �C.

Figure 6 XRD peaks of Al ®lms deposited at various substrate

temperatures.

287

XRD spectra of the Al ®lms deposited at various

substrate temperatures on TiN. The ratio of

I�1 1 1�=I�2 0 0� was between 2.34 and 3.04. The Al thin

®lms showed a (1 1 1) preferred orientation at all growth

temperatures. It is well known that Al (1 1 1) ®lm

formation was induced by the (1 1 1) oriented CVD-TiN

underlayer. In this study, Al ®lms were deposited on

(2 0 0) oriented TiN ®lms and ®lms did not show the high

value of I�1 1 1�=I�2 0 0�. The intensity ratio did not change

signi®cantly with the growth temperature.

In order to study the morphology of the deposited Al,

scanning electron microscopy (SEM) and atomic force

microscopy (AFM) were used. The SEM images of

deposited Al ®lms at 140 �C and 190 �C with a thickness

of 900 nm are shown in Fig. 7. The microstructure of Al

®lms was signi®cantly changed as a function of the

growth temperature. Generally, changes in grain growth

with increasing temperature arise from thermally

activated processes. At 140 �C, small and ®ne crystallites

formed and these closely packed grains act as a

continuous layer on which further nucleation of Al can

occur. Above 140 �C, grains of Al increase in size and

begin to coalesce as the substrate temperature is

increased. However, the grain size did not change

much at 160±190 �C due to the fact that a higher

temperature is needed for the activation of Al recrys-

tallization and signi®cant grain growth.

The root mean square (rms) value of surface roughness

as a function of the substrate temperature was calculated

from AFM measurements of Al ®lms with a thickness of

Figure 7 SEM photograph of CVD-Al ®lms deposited at (a) 140 �C and (b) 190 �C.

288

about 900 nm and is plotted in Fig. 8. Below 190 �C, the

rms surface roughness was not changed signi®cantly but

surface roughness increased linearly with the increase of

the substrate temperature above 190 �C.

Fig. 9 shows the effect of the ®lm thickness of CVD Al

®lms deposited at 190 �C on ®lm roughness. It is obvious

that topology becomes worse as the ®lm thickness

increases, probably due to the island growth character-

istics. During the early stage of growth, the surface

roughness increased rapidly with increase in ®lm

thickness, while after the thickness reached * 300 nm,

the surface roughness increased gradually due to the

disappearance of the TiN substrate effect.

4. ConclusionsWe synthesized TMEDAA by the ligand displacement

method and the metalorganic chemical vapor deposition

of Al ®lms from TMEDAA was conducted to examine

the deposition rate, electrical resistivity, microstructure,

surface roughness, and the preferred orientation. The

molecular structure of TMEDAA is polymeric in the

solid state but cyclic dimeric in the gas phase. In the

temperature range 140±210 �C, the Arrhenius plot

showed a reaction-rate-limited regime with an activation

energy of 58.6 kJ/mol. At high temperatures, incorpora-

tion of impurities increased, resulting in higher

resistivity. There was a signi®cant change in the ®lm

morphology as a function of the substrate temperature

and ®lm thickness. As the ®lm thickness increases, the

surface roughness increased and as the substrate

temperature was increased, a small and closely packed

grain structure turned into a faceted grain structure along

with an increase of surface roughness. The Al ®lms

showed a (1 1 1) preferred orientation at all substrate

temperature.

Figure 8 Rms (root mean square) roughness of Al ®lms with thickness

of 900 nm deposited at various substrate temperatures.

Figure 9 AFM images of Al ®lms (a) 80 nm, (b) 220 nm, (c) 900 nm, and (d) rms (root mean square) roughness as a function of ®lm thickness.

289

AcknowledgementsThe authors would like to thank LG Semicon Co., Ltd.

and Korea Research Foundation (KRF) for their support

of this study. The authors are grateful to Dr J. W. Park and

C. H. Lee for making the aluminum precursor.

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Received 10 August 1998and accepted 14 December 1998

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