The effects of substrate heating and substrate biasing on the initial stage of nonepitaxial heterogeneous growth of TiN on Si(111) was studied by using high-resolution transmission electron microscopy. Although TiN films deposited at room temperature (RT) undergo a transition from continuous amorphous films to polycrystalline films with three-dimensional grains when the film thickness is increased from ~1 to 2 nm, crystallization occurred at a substrate temperature, T/sub s/=570 K, even for film thicknesses less than 1 nm. Compared with growth at T/sub s/=RT, at T/sub s/=570 K, the initial lateral grain size was only slightly larger, and the grains tended to be spherical and discontinuous at higher film thickness. At a substrate bias voltage, V/sub b/=-70 V, the grains were laterally larger and planar. At a film thickness of 50 nm, the films deposited at V/sub b/=-70 V showed the thermodynamically favored (200) preferred orientation, whereas the films deposited at T/sub s/=570 K showed (111) preferred orientation with a weak (200) peak

Komiyama Hiroshi

Li Tu-Qiang

Noda Suguru

Okada Fumio

小宮山 宏

岡田 文雄

野田 優

UTokyo Repository

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Effects of substrate heating and biasing on nanostructural 
evolution of non-epitaxially grown TiN nano-films 
 
T. Q. Li a), S. Noda a), F. Okada, and H. Komiyama 
 
Department of Chemical System Engineering, University of Tokyo, 7-3-1 Hongo, 
Bunkyo-ku, Tokyo 113-8656, JAPAN 
 
(Received 
                                                  
a)Author to whom correspondence should be addressed; electronic mail: 
litg@chemsys.t.u-tokyo.ac.jp; noda@chemsys.t.u-tokyo.ac.jp 
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Abstract 
The effects of substrate heating and substrate biasing on the initial stage of 
non-epitaxial heterogeneous growth of TiN on Si(111) was studied by using 
high-resolution transmission electron microscopy. Although TiN films 
deposited at room temperature (RT) undergo a transition from continuous, 
amorphous films to polycrystalline films with three-dimensional grains 
when the film thickness is increased from ~ 1 to 2 nm, crystallization 
occurred at a substrate temperature, Ts = 570 K, even for film thicknesses 
less than 1 nm. Compared with growth at Ts = RT, at Ts = 570 K the initial, 
lateral grain size was only slightly larger, and the grains tended to be 
spherical and discontinuous at higher film thickness. At a substrate bias 
voltage, Vb = -70 V, the grains were laterally larger and planar. At a film 
thickness of 50 nm, the films deposited at Vb = -70 V showed the 
thermodynamically favored (200) preferred orientation, whereas the films 
deposited at Ts = 570 K showed (111) preferred orientation with a weak 
(200) peak. 
 
PACS numbers:  
1. 68.55.Ac. 
2. 68.55.Jk. 
3. 81.15.Cd. 
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I. INTRODUCTION 
Due to size miniaturization and density increase of integrated circuits (ICs), the 
thickness of the barrier layers (TiN, TaN, etc.) for metal interconnections (Al, Cu, etc.) 
must be reduced to less than 10 nm in order to reduce the resistivity of metal/barrier 
structures.1 To develop such ultra-thin diffusion barriers, a detailed understanding of 
their initial growth mechanism is needed. 
The barrier layers are often deposited non-epitaxially on either Si or SiO2 substrates 
and form through crystal nucleation, grain growth, and coalescence. Many 
investigations have been made on the texture evolution2-4 and homogeneous growth of 
TiN.5-7 Unfortunately, little effort has been directed at understanding initial formation 
mechanisms of these non-epitaxial heterostructures.8,9 Previous results indicated that 
TiN deposited onto Si by reactive sputtering undergoes a thickness-dependent crystal 
nucleation at room temperature.9 Further studies are needed, therefore, to determine the 
effect of deposition parameters on the nonepitaxial, initial growth. For thick film growth 
(e.g., at micrometer-ordered thickness), substrate heating or biasing are known to favor 
the formation of (200) preferred orientation because this is the lowest surface energy 
existing for the (200) plane of NaCl-type crystal TiN films.10 However, at the initial 
growth stage when the films have a thickness on the order of nanometers, the effects are 
more complicated due to non-epitaxtial heterogeneous nucleation and coalescence. In 
this work, the influences of substrate temperature, Ts, and bias voltage, Vb, on the initial 
film nanostructure are investigated by using high-resolution transmission electron 
microscopy (HRTEM). 
 
II. EXPRIMENTS 
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TiN thin films were deposited using dc reactive magnetron sputtering. The sputter 
chamber had a base pressure of 5×10-6 Pa. The target was a 2-inch-diameter, 99.99% 
pure Ti plate, and the substrates on which the films were grown were (111) Si wafers. 
Before being loaded into the loading chamber, the wafers were treated chemically in a 
mixture of concentrated sulfuric acid and H2O2, and then dipped into a 1% HF solution 
to remove the contaminants and the native SiO2 layer on the surface. After each 
chemical treatment, the wafers were washed in deionized water. Before deposition, the 
target was pre-sputtered for 5 min under the same conditions as for the TiN deposition. 
The substrate was shuttered with a rotating shuttering system during pre-sputtering. 
Deposition was done by rotating the shutter to a position where a hole in the shutter was 
in line between the target and the substrate. During deposition, the total pressure, P, in 
the sputtering chamber was maintained at P = 0.93 Pa, and the N2 partial pressure, PN2 = 
0.047 Pa. The substrate-to-target distance was fixed at 50 mm, and the discharge power 
to the Ti target was 69 W dc (260 V, 0.265 A). TiN films were deposited under these 
conditions without substrate heating (in this paper we refer to this condition as Ts = RT, 
where RT represents room temperature and 290 K < RT < 308 K), or at Ts = 570 K. The 
substrate bias voltage, Vb, was either a self-bias voltage, Vb = - 23 V, induced from the 
plasma, or was set to Vb = - 70 V. The ion current density at the substrate, determined 
from the current density-voltage characteristic of the substrate (not shown here), was 
estimated to be 0.57 mAcm-2 (corresponding to an ion flux of 3.5×1015 cm-2s-1). 
The nanostructure of the films was investigated with transmission electron 
microscopy (TEM: JEOL, JEM2010F, Japan) operated at 200 kV. HRTEM observations 
of plan-view and cross-sectional (XHRTEM) specimens were made. 
The deposition rate was determined by measuring the film thickness (about 50 nm) 
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with a stylus profilometer (KLA-Tencor P-10, U.S.A.). The thickness of initial films, hc, 
was calculated from the deposition rate (R) and time (t) as hc = R × t. The deposition 
rates at Vb = - 23 V and -70 V were 0.19 and 0.17 nm/s, respectively. The slightly lower 
deposition rate at Vb = - 70 V was presumably due to resputtering on the film surface. 
The x-ray diffraction (XRD) (Rigaku RINT2400, Japan) analysis was done by using 
Cu Kα radiation. 
Additional details of our experimental setup are given elsewhere.8, 9 
 
III. RESULTS 
A. Growth at 570 K 
Figs. 1 and 2 (modify the arrows ?!) show plan-view and cross-sectional HRTEM 
images of films deposited at Ts = 570 K (Vb = -23 V). Figs. 1(a) and 2(a) show films 
deposited for t = 4 s (hc = 0.8 nm). Samples deposited at RT were continuous and 
amorphous, as determined by HRTEM.9 However, samples deposited at Ts = 570 K 
exhibited crystal nuclei in the sub-nanometer-film. The XHRTEM images show that the 
nuclei were nearly spherical and that their height was about 2 nm, more than twice that 
of hc. This implies that deposits aggregated at Ts = 570 K. 
For t = 8.5 s (hc = 1.6 nm), Figs. 1(b) and 2(b) show that the nuclei (grains) grew 
three dimensionally, resulting in a rougher film surface compared with films deposited 
at Ts = RT. (see Fig. 2 in Ref. 9) 
Figs. 1(c) and 2(c) show that for t = 25 s (hc = 4.8 nm), the grains grew large to come 
in contact with each other, resulting into a polycrystalline film with distinct grain 
boundaries. The SAED patterns indicate that the (do not use so many terms such as 
“nuclei”, “grains”, and “crystallites”) grains had random orientation both in-plane and 
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out-of-plane directions. 
For hc = 16 nm, the XRD results shown in Fig. 3(b) indicate that the film had a nearly 
random orientation and that the (200) intensity was stronger than that of the film 
deposited at Ts = RT [shown in Fig. 3(a)]. The intensity of the (200) peak remained 
unchanged with increasing hc, whereas the intensity of the (111) peak increased rapidly 
with increasing hc, and the film exhibited a (111) preferred orientation at hc = 40 nm. 
For Ts = 570 K, the XHRTEM images show an amorphous interlayer 0.9 nm thick 
between the deposits and the substrate, whereas for films deposited at Ts = RT, the 
interlayer thickness was less than 0.5 nm thick (see Fig. 2 in Ref. 9). XHRTEM and 
XPS analysis showed that the interlayer formed at Ts = RT and PN2 = 0.47 Pa was 
mainly SiNx.9 The formation of the thicker amorphous interlayer at Ts = 570 K may be 
due to enhanced inter-diffusion and/or oxidation of the Si surface when heated in the 
sputter-chamber before deposition of TiN. Fig. 2 indicates that there should be little Ti 
in the interlayer, because the contrast of the interlayer was much brighter than that of 
the TiN layer, and closer to that of Si. If the interlayer contained high concentrations of 
Ti, the relatively heavy Ti atoms would efficiently scatter electrons, and therefore 
darken the contrast of the layer.  
Fig. 4 shows a correlation of the changes of the average lateral grain size with film 
growth at Ts = RT and Ts = 570 K. Each reported lateral grain size is an average of the 
measured size of more than 100 randomly-sampled grains from plan-view HRTEM 
images taken from a single film sample. The error bars represent the standard deviation 
of the lateral grain size. For Ts = 570 K, from hc = 0.8 to 1.6 nm the lateral grain size 
increased from 2.2 to 3.0 nm, and the lateral grain size was almost the same as the grain 
height. This indicates three-dimensional grain growth and is identical to the XHRTEM 
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results shown in Fig. 2. At hc = 4.8 nm, the lateral grain growth rate decreased because 
the grains came into contact with each other, and mainly grew vertically. 
Fig. 4 indicates that compared with the grains deposited at Ts = 570 K, the grains 
grown at Ts = RT were smaller and had lower lateral growth rate. This may be because 
of the lower adatom mobility and higher density of initial nucleation sites at the lower 
substrate temperature. The grains therefore came into contact with each other earlier in 
the film-growth process. 
 
B. Growth under substrate biasing 
Fig. 5 show results for TiN initial films deposited at Vb = -70 V (Ts = RT). Fig. 5(a) 
shows a plan-view HRTEM image together with SAED patterns for a film with hc = 3.4 
nm. The grains had an average lateral size of 5.9 nm, about twice the film thickness. As 
shown in Fig. 5(a), some grains with relatively large lateral size were (200) oriented. 
Fig. 5(b) shows an XHRTEM image for a film with hc = 2.5 nm. The film was smooth 
and the grains were planar with a relatively large top surface area. 
Fig. 6 shows the XRD spectra for hc = 50 nm, and for Vb = -23 V (self-bias potential) 
and Vb = -70 V. For Vb = -23 and -70 V, the film orientation was (200) and (111), 
respectively. 
 
IV. DISCUSSION 
Previously reported HRTEM results showed that at Ts = RT (Vb = -23 V), the initial 
non-epitaxial growth of TiN on Si undergoes a thickness-dependent 3D crystal 
nucleation from an initial continuous, 2D amorphous layer.9 For Ts = 570 K, 
crystallization occurred for hc < 1 nm. This may be because with increasing Ts, the 
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adatom mobility increases, causing the formation of islands. With increasing Ts, the 
amount of film material required for the discontinuous islands to reach a critical height 
decreases, permitting crystal nucleation to occur in the initial stages of deposition. 
The texture of the initial layers is affected by two important energy parameters: the 
film/substrate interfacial energy, γif/s, and the film surface energy, γsf. The minimization 
of both γif/s and γsf is the primary driver for texture development in the initial stage. For 
epitaxial growth, lattice matching causes a large decrease in γif/s, and the deposited film 
orientation is determined by matching the orientation of the substrate. For non-epitaxial, 
heterogeneous growth, the initial film texture is mainly determined by minimization of 
γsf. In the initial growth stage when the crystal nuclei (grains) are discontinuous, γsf does 
not depend on the grain orientation and no orientation is preferred if the grains are 
non-wetting and spherical. This is consistent with previous results9 and with the results 
shown in Section III that TiN growth at either Ts = RT or 570 K shows no preferred 
orientation in the nucleation stage. 
When the crystal grains grow large enough to contact each other, the continuous film 
tends to orient itself with the lowest surface-energy plane parallel to the substrate to 
minimize γsf. At Ts = RT and low ion irradiation, this thermodynamically favored 
transition does not occur to any significant degree, and the grains mainly grow 
vertically with an initial, random orientation. Because the (111) oriented grains have a 
higher growth rate than the other oriented grains, (111) oriented films are formed by the 
“evolutionary selection rule”. 4, 8, 11 
When the films are deposited at Ts = 570 K, the grains tend to become spherical and 
discontinuous. In addition, the grains become continuous at a greater thickness than for 
Ts = RT. The evolutionary selection growth, which leads to (111) preferred orientation, 
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therefore occurs later than for Ts = RT. This delay in the occurrence of evolutionary 
selection growth might be the cause of the stronger intensity of the (200) orientation in 
the XRD spectra of films deposited at Ts = 570 K than for films deposited at Ts = RT. 
Fig. 5 indicates that grains of the films deposited at Vb = -70 V have a larger lateral 
than vertical size, and are planar. For films deposited with substrate biasing the ions 
(such as Ar+ and N2+, etc.) impinge on the film surface with higher kinetic energy. This 
higher ion energy at Vb = -70 V may cause the following effects on the film, resulting in 
the formation of planar grains. Higher ion energy: 
(1) increases the mobility of the adatoms on the grain surface, thus promoting 
diffusion of atoms to the edges.  
(2) promotes grain coalesce with adjacent grains, forming a planar shape. 
(3) causes loosely bound ions on the surface of the grains to be resputtered, 
resulting in a denser film structure. 
Due to their larger top than side area, to lower their surface energy planar grains may 
restructure into (200) orientation planes under energetic ion bombardment. The films 
show (200) preferred orientation at a thickness of 50 nm. 
 
V. CONCLUSIONS 
Although TiN films deposited at room temperature (Ts = RT) undergo a transition 
from continuous, amorphous layers to three-dimensional, crystalline layers, when the 
film thickness increased from ~ 1 to 2 nm, crystallization occurred at Ts = 570 K, even 
for film thicknesses less than 1 nm. Compared with growth at Ts = RT, at Ts = 570 K the 
initial grains tended to be spherical and discontinuous at higher film thickness. At a 
substrate bias voltage, Vb = -70 V, the grains were laterally larger and planar. The 
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different effects of Ts and Vb on TiN initial growth resulted in different textures at a film 
thickness of 50 nm 
 
ACKNOWLEDGEMENTS  
This work was supported, in part, by the New Energy and Industrial Technology 
Development Organization’s (NEDO) "Nanotechnology Program - Systemization of 
Nanotechnology Materials Program Results Project", of the Ministry of Economy, Trade, 
and Industry (METI), Japan, and by the Japan Society for the Promotion of Science 
(JSPS)’s "Grant-in-Aid for Creative Scientific Research - Establishment of Networked 
Knowledge System with Structured Knowledge for Future Scientific Frontier Project", 
of the Ministry of Education, Culture, Sports, Science and Technology (MEXT), Japan. 
 
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1S. P. Murarka, R. J. Gutmann, A. E. Kaloyeros, and W. A. Lanford, Thin Solid Films 
236, 257 (1993). 
2U. C. Oh and J. H. Je, J. Appl. Phys. 74, 1692 (1993). 
3J. Pelleg, L. Z. Zevin, and S. Lungo, Thin Solid Films 197, 117 (1991). 
4L. Hultman, J.-E. Sundgren, J. E. Greene, D. B. Bergstrom, and I. Petrov, J. Appl. Phys. 
78, 5395 (1995). 
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S. Kodambaka, P.O’Sullivan, and I. Petrov, MRS Bull. 26, 182 (2001). 
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Greene, Thin Solid Films 392, 164 (2001). 
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Rev. Lett. 89, 176102-1 (2002). 
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Lett. 67, 2928 (1996). 
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FIGURE CAPTIONS 
Fig. 1 Plan-view TEM images and corresponding SAED patterns of TiN deposited at a 
temperature of 570 K for deposited film thicknesses of (a) 0.8 nm, (b) 1.6 nm, and (c) 
4.8 nm. 
Fig. 2 XHRTEM images of TiN deposited at a temperature of 570 K for deposited film 
thicknesses of (a) 0.8 nm, (b) 1.6 nm, and (c) 4.8 nm. 
Fig. 3 XRD scans showing the effect of film thickness on preferred orientation of TiN 
deposited at (a) room temperature and (b) 570 K. 
Fig. 4 Lateral grain size vs. film thickness. 
Fig. 5 (a) Plan-view TEM image and corresponding SAED patterns of a 3.4-nm film 
and (b) XHRTEM image of a 2.5-nm film deposited at a bias voltage of Vb = -70 V. 
Fig. 6 Effect of substrate bias voltage on the texture of 50-nm films. 
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Fig. 1 
 
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Fig. 2 
 
 
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Fig. 3 
 
 
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Fig. 4 
 
 
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Fig. 5 
 
 
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Fig. 6 
 
 
 


Effects of substrate heating and biasing on nanostructural evolution of nonepitaxially grown TiN nanofilms

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Effects of substrate heating and biasing on nanostructural evolution of nonepitaxially grown TiN nanofilms

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