Full text of DOI:10.1063/1.115268

Structural characterization of aluminum films deposited on sputteredtitanium
nitride/silicon substrate by metalorganic chemical vapor deposition from
dimethylethylamine alane
Xiaodong Li, ByoungYoup Kim, and ShiWoo Rhee
Citation: Applied Physics Letters 67, 3426 (1995); doi: 10.1063/1.115268
View online: http://dx.doi.org/10.1063/1.115268
View Table of Contents: http://scitation.aip.org/content/aip/journal/apl/67/23?ver=pdfcov
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Structural characterization of aluminum films deposited
on sputtered-titanium nitride/silicon substrate by metalorganic
chemical vapor deposition from dimethylethylamine alane
Xiaodong Li, Byoung-Youp Kim, and Shi-Woo Rheea)
Laboratory for Advanced Materials Processing (LAMP), Department of Chemical Engineering,
Pohang University of Science and Technology (POSTECH), San 31 Hyoja Dong, Pohang,
Kyung-buk 790-784, Republic of Korea
͑
Received 31 July 1995; accepted for publication 27 September 1995͒
Al films deposited on sputtered-TiN/Si substrate by metalorganic chemical vapor deposition
͑
͑
MOCVD͒ from dimethylethylamine alane ͑DMEAA͒ were characterized using x-ray diffraction
XRD͒, Auger electron spectroscopy ͑AES͒, atomic force microscopy ͑AFM͒, and transmission
electron microscopy ͑TEM͒. The TiN film sputtered on the Si has a preferred orientation along the
growth direction with the ͗111͘ of the film parallel to the Si͗111͘. Sputtering of the TiN film on the
Si induced strains at the interface. The TiN/Si interface is flat while the Al/TiN interface is rough.
There exist many dislocations at the Al/TiN interface. The Al O phase was formed at the Al/TiN
2
3
interface during the early stages of Al deposition. In the Al grains, there exist many tangled
dislocations and a few Al O particles. With increasing deposition time, the Al film surface
2
3
roughness increases. © 1995 American Institute of Physics.
Multilevel metallization is an indispensable technique in
increasing the device packing density in very large-scale in-
tegrated circuits ͑VLSIs͒. Al and Al alloy films are widely
used for such metallization. Generally, Al is deposited by
physical vapor deposition ͑PVD͒, however, as aspect ratio
increases, it is difficult for PVD to fill via holes of deep-
submicron size. Chemical vapor deposition ͑CVD͒ of Al has
the advantages of excellent step coverage and selective
growth, which are convenient for such micrometal-
reactor with vapor phase mass flow controller without any
Ϫ6
carrier gas. The base pressure was 1.1ϫ10
Torr and pro-
cess pressure was 0.15 Torr.
The structures of the TiN and Al films were analyzed
using a Rigaku D-Max 1400 x-ray diffractometer with
Cu K␣ radiation operating at 50 kV, 150 mA. The depth pro-
files of elements Al, O, N, Si, and Ti were obtained by a
Perkin-Elmer PHI 600 scanning Auger multiprobe. The sur-
face morphology of the TiN and Al films was observed using
1
–4
lization.
DMEAA is one of the promising candidates for Al CVD
because of its relatively high vapor pressure at room tem-
perature ͑1.5 Torr͒, its long shelf-life, and its being a liquid
and thus providing stable vapor pressure. Moreover, direct
Al–C bonds are not present in its molecular structure, which
3
helps to prevent carbon contamination. To improve contact
properties, a barrier layer is required for Al metallization in
VLSIs. TiN is frequently used as a barrier layer because of
its thermal and chemical stability. Although Al CVD on
5
metal films using DMEAA has been reported, the structure
of the Al films is still far from clear. We deposited TiN film
on Si as a diffusion barrier with reactive sputtering, and then
deposited Al films on the TiN/Si substrate by CVD with
DMEAA.
The Si͑100͒ wafers having 4° off angle to ͑100͒ were
used in this study. After a standard cleaning process, TiN film
of 500 nm thickness was formed on the Si by direct current
magnetron reactive sputtering in a gas mixture of Ar and
N . Sputtering parameters were as follows: base pressure 5
2
Ϫ9
ϫ10
Torr, process pressure 3 mTorr, sputtering power 6
kW, N flow rate 60 sccm, Ar flow rate 20 sccm, substrate
2
heated to 400 °C during sputtering. Al films were prepared at
190 °C by CVD with DMEAA. DMEAA was fed into the
FIG. 1. XRD spectra of the Al films deposited on the TiN/Si substrate for ͑a͒
15 and ͑b͒ 1 min, respectively.
a͒Electronic mail: srhee@vision.postech.ac.kr
3
426
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Appl. Phys. Lett. 67 (23), 4 December 1995
0003-6951/95/67(23)/3426/3/$6.00
© 1995 American Institute of Physics
1

FIG. 2. Typical AES depth profiles from an Al film deposited on the TiN/Si
substrate.
an Autoprobe-CP model atomic force microscope. Cross-
sectional specimens for TEM were first glued face to face,
cut, then ground, dimpled, and thinned by ion milling. TEM
observations were conducted using a JEOL JEM-200CX
transmission electron microscope operating at 100 kV.
Figure 1 shows the XRD spectra of the Al films depos-
ited on the TiN/Si substrate for 1 and 15 min, respectively.
The thickness of the Al films deposited on the TiN/Si sub-
strate for 1 and 15 min was estimated to be about 50 and
1000 nm, respectively, by weight measurement. The TiN film
sputtered on Si exhibits strong ͑111͒ orientation while the Al
films deposited on the TiN/Si substrate show random orien-
tation. This is not in agreement with the results of Al films
sputtered on the sputtered-TiN͑111͒ films, where the Al films
6
give ͑111͒ orientation. The Al film deposited on the TiN/Si
substrate for 1 min shows several peaks from Al O in addi-
2
3
tion to the peaks of both Al and TiN while the Al film de-
posited on the TiN/Si substrate for 15 min does not. The
Al O phase is not likely to be from surface native oxide
2
3
because it was not detected by XRD when the film was thick.
This indicates that the Al O phase or intermediate layer
2
3
formed during the early stages of Al deposition.
Figure 2 shows typical AES depth profiles from an Al
film deposited on the TiN/Si substrate. A small oxygen peak
was observed at the Al/TiN interface near the Al film side.
This suggests that Al–O compounds formed at the Al/TiN
interface, in good agreement with the XRD results. A Ti–O
compound layer was reported to exist on the TiN film
FIG. 3. AFM images of ͑a͒ the TiN film, and the Al films deposited on the
TiN/Si substrate for ͑b͒ 15 and ͑c͒ 1 min, respectively..
7
surface. The formation of the Al O phase at the Al/TiN
2
3
interface suggests that during Al CVD Ti–O compounds re-
acted with Al, forming the Al–O compounds. The effect of
temperature and other operating variables on the formation
of alumina at the interface is not yet known.
Figure 3 shows the AFM images of the TiN film, and the
Al films deposited on the TiN/Si substrate for 1 and 15 min,
respectively. The TiN and Al films exhibit rugged surface
morphology. The rugged mountains of the TiN film distribute
uniformly while those of both Al films do not. With increas-
ing deposition time, the Al film surface roughness increases.
The average diameters of the rugged mountains of the TiN
film, and the Al films deposited on the TiN/Si substrate for 1
and 15 min, are 1020, 2160, and 5630 nm, respectively.
From the AFM results, however, it is difficult to determine
the structure of the TiN and Al films.
Figure 4 shows the bright field cross-sectional TEM im-
ages of the Al film deposited on the TiN/Si substrate for 15
min and the Al/TiN interface, respectively. Good sections
with homogenous thickness were obtained with the TiN and
Al films ͓Fig. 4͑a͔͒. The thickness of the Al film was found
to be about 1000 nm, and that of the TiN film was found to
be about 500 nm. The TiN film was found to be made up of
very small needlelike grains ϳ20 nm in width and 200 nm in
length, which grow in the direction inclined to the Si at an
Appl. Phys. Lett., Vol. 67, No. 23, 4 December 1995
Li, Kim, Rhee
3427
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1

found to have elongated blocklike shape ϳ600 nm in width
and 930 nm in length. Some small Al grains, which were
formed at the nucleation stage, were found near the interface.
In the Al grains, there exist many tangled dislocations, and a
few small particles with a diameter of about 50 nm which
were identified as Al O phase by selected area diffraction
2
3
analysis. Al O particles cannot nucleate and grow to such
2
3
size in the Al grains within 30 min at 190 °C. It was also
found that there is not any definite orientational relationship
between the Al O particles and Al matrix. Therefore, a pos-
2
3
sible reason for the existence of Al O particles within the Al
2
3
grains is that Al O particles formed in gas phase because of
2
3
residual O or H O and then deposited together with Al onto
2
2
specimen surface. DMEAA is unstable in the atmosphere and
it readily decomposes into alumina when exposed to air. Be-
cause the volume fraction of Al O particles is small, it is
2
3
impossible to detect them by XRD. The existence of Al O3
2
particles in the Al grains is also one of the reasons why there
exist many tangled dislocations in the grains. Compared with
the AFM results, it was found that the rugged mountains in
the AFM images do not correspond to the grains. It is desir-
able that grain size should be large and uniformly distributed.
Also, the film should be free from impurities. Other research-
8
9
ers showed that impurities like hydrogen, nitrogen, and
oxygen10 have bad effects on the migration resistance of
metal lines.
In summary, the TiN film sputtered on the Si has a pre-
ferred orientation along the growth direction with the ͗111͘
of the film parallel to the Si͗111͘. Sputtering of the TiN film
on the Si induced strains at the interface. The TiN/Si inter-
face is flat while the Al/TiN interface is rough. There exist
many dislocations at the Al/TiN interface. Al O3 phase
2
formed at the Al/TiN interface during the early stages of Al
deposition. In the Al grains, there exist many tangled dislo-
cations and a few Al O particles. With increasing deposition
2
3
time, the Al film surface roughness increases.
This research was supported by LG Semiconductor Co.
and support of Engineering Research Center for Interface
Science and Technology of Materials is also acknowledged.
FIG. 4. Bright field cross-sectional TEM images of ͑a͒ the Al film deposited
on the TiN/Si substrate for 15 min and ͑b͒ the Al/TiN interface, respectively.
1
N. Takeyasu, Y. Kawano, E. Kondoh, T. Katagiri, H. Yamamoto, H. Shin-
angle of about 66°. This direction is parallel to the Si͗111͘.
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tion of the TiN film is ͗111͘. Therefore, there is a preferred
orientation of the TiN film along the growth direction with
the film ͗111͘ parallel to the Si͗111͘. This is probably be-
cause atomic arrangement and interatomic distance of the
TiN͑111͒ plane are close to those of the Si͑111͒. Obvious
changes in contrast were observed in the Si near the interface
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relatively rough, and there exist many dislocations ͓Fig.
8
9
4
4
͑b͔͒. An intermediate layer indicated by the arrow in Fig.
͑b͒ seems to exist at the Al/TiN interface, which might be
10
the Al O phase found by XRD and AES. The Al grains were
2
3
3428
Appl. Phys. Lett., Vol. 67, No. 23, 4 December 1995
Li, Kim, Rhee
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