Critical metrology for ultrathin high k dielectrics.
W.Vandervorsta'b, B.Brijs, H.Bender, T.Conard, J.Petry, O.Richard, X.
Blascoc, M. Nafriac
a
IMEC, Kapeldreef 75, B-3001 Leuven, Belgium
b
also : INSYS, KU Leuven, Belgium
c
Departament d'Enginyeria Electronica, Universitat Autonoma de Barcelona,
Edifici Q. 08193. Bellaterra. Spain
Abstract. Targeting very thin equivalent oxides (<1 nm) requires the deposition of (very) thin dielectrics onto
silicon surfaces with minimal interfacial oxide. Typically, high-k dielectric layers are deposited using ALD or MOCVD
with, at present, a prime emphasis on Hf-based high-k dielectrics, either as pure HfO2, as silicate, or mixed with A12O3.
Depending on the deposition conditions, serious deficiencies in terms of film closure and material density occur for the
ultra thin (<3 nm) films which are required for very small EOT's. As such, critical metrology needs arise enabling one
to study details of the film growth and its evolution upon thermal anneal. As discussed in this paper, many tools are
required such as Rutherford Backscattering Spectrometry and high-resolution elastic recoil detection (ERDA), Low
Energy Ion Scattering, Time-of-flight SIMS, (spectroscopic) ellipsometry and X-ray photoelectron spectroscopy. When
trying to correlate these different methods one must be aware of potential discrepancies due to non-homogeneous
growth and reduced material density. Local electrical measurements based on tunneling atomic force microscopy reveal
very fine scale inhomogeneities, which can be correlated to local structural defects.
MOCVD HfO2 and Hf-silicates are deposited
respectively
using tetrakis
(diethylamino)
hafnium (TDEAH) and TDMAS tetrakis
(dimethylamino) silicon) as precursor. In all
cases it is clear that achieving a high quality thin
film requires a complete control over the growth
process and the means to actually monitor the
film growth and its structural and electrical
properties. A complete review of all possible
characterization methods of value in this context
is beyond the scope of this paper. Instead we will
focus on the issues of determining film thickness
and growth mode, film composition and density,
and local electrical properties.
INTRODUCTION
High-k insulating metal oxides are currently
under consideration as novel gate dielectric
materials for silicon MOS devices because they
provide the required specific capacitance at a
considerably larger physical thickness than SiO2,
thus allowing the reduction of the gate leakage
current by suppression of direct tunneling.
Precise control of the layer growth, the evolution
of the interface between the silicon and the highk dielectric and the thermal stability of the gate
stack are key elements towards meeting the
target of the equivalent oxide thickness being
less than 0.5 nm for very advanced applications.
Even when using high k materials such as A12O3,
ZrO2, HfO2, and their mixtures with Si or Al
which have a value of k ranging from 9-25,
obtaining an EOT of 0.5 nm still requires very
thin films i.e. in the order of 2-3 nm ideally with
no, or minimal, interfacial oxide.
STRUCTURAL FILM METROLOGY
Film thickness and density
One of the obvious approaches to study the basic
properties of these deposition processes is to
look at growth curves i.e. the evolution of film
thickness as a function of growth/deposition
time. The most unambiguous way is to rely on
techniques like Rutherford Backscattering (RBS)
or X-ray Fluorescence (XRF) which measure
directly the number of atoms deposited per cm2.
Typical examples of growth curves (for ALD
deposition of A12O3, HfO2 or ZrO2 on an HF-last
surface) are shown in Fig. 1. They clearly
indicate that film growth can occur in a very
The depositions of high k films, as discussed in
this work, are either based on Atomic Layer
Deposition (ALD) or metal-organic chemical
vapor deposition (MOCVD). In its most
simplistic way, ALD can be described as a series
of sequential saturating gas-solid reactions by
alternating H2O and precursor pulses (in this
work (A1(CH3)3, ZrCl4 or HfCl4) which should
result in excellent uniformity of the deposited
layers and accurate film thickness control. In
Contacting author : [email protected]
CP683, Characterization and Metrology for VLSI Technology: 2003 International Conference,
edited by D. G. Seiler, A. C. Diebold, T. J. Shaffner, R. McDonald, S. Zollner, R. P. Khosla, and E. M. Secula
© 2003 American Institute of Physics 0-7354-0152-7/03/$20.00
129
non-linear mode and apparently is very sensitive
to the surface preparation. With such a nonlinear growth behavior it is clear that tuning of
the deposition systems and recipes becomes
important and thickness control will be a key
activity in a production line and during process
development. For that purpose a fast noncontacting technique like ellipsometry would be
a preferred choice above (expensive) large
accelerators, in particular since ellipsometry is,
with success, already extensively used in the
normal semiconductor processing environment
[1]. As the gate stack is typically a multilayer
structure (some interfacial oxide is almost
always present) multilayer fitting is a necessity
[2]. Despite the use of spectroscopic ellipsometry
and extending the wavelengths to the vacuum
UV range (135- 620 nm range), a large
uncertainty on the interfacial oxide thickness
remains and its value usually needs to be fixed in
order to obtain a reliable value for the high k
film thickness [3].
t =
thickness t [5] (eq.l), the angular range needs to
be limited (depending on film thickness) to less
than 60 degrees as deviations from the expected
linear relationship with cos(0) are observed for
larger angles (probably due to elastic scattering)
(Fig. 2) [6].
9,0 nm
1.9 nm
3
4
1/008(9)
0
45
60
70
eo
75
Figure 2. AR-XPS-results obtained on HFO2-films on
Si. (Data are collected with a ThermoVG Scientific
Theta probe) [4]. The deviation from linearity for large
angles indicates the breakdown of eq.(l) for those
angles. [6]
o
50
When trying to establish a correlation between
techniques like (spectroscopic) ellipsometry
(SE), X-Ray Photoelectron Spectroscopy (XPS),
or Transmission Electron Microscopy (TEM),
i.e. between techniques measuring directly film
thickness as a length scale (SE, TEM,..), versus
techniques measuring the number of atoms/cm2
(RBS), the uncertainty on the density of the very
thin films becomes an essential parameter. In this
context it is also important to emphasize that the
thickness as determined by XPS based on the
attenuation of the Si-substrate signal, inherently
contains an assumption on the density as long as
standard attenuation lengths are used in the
calculations [7].
In principle X-ray reflectometry (XRR) allows
the simultaneous determination of thickness,
density and roughness of (multi-)layers by fitting
of the measured spectra with calculated ones.
Unfortunately for very thin layers (< 1-2 nm) the
XRR spectrum shows little structure and the
correlation between the different parameters
becomes large, i.e. the accuracy of the fitting
might be questionable. Nevertheless, when
comparing the XRR-results with the density
extracted from the number of atoms as measured
10
number of cycles
Figure 1. Growth curve measured by RBS for A12O3,
ZrO2 and HfO2 grown by ALD on a HF- last Si or
SiO2-surface.
An inherent limitation with the application of the
fitting models is that use is made of optical data
extracted from thicker films, which may or may
be not be applicable to very thin films.
In principle an alternative approach would be to
use angular resolved X-Ray Photoelectron
Spectroscopy (XPS) providing information on
the high k film thickness as well as on the
interfacial oxide. The latter can presently be
done on a full wafer scale either by repeating the
measurements with different acceptance angles
or by using instruments which provide directly
multi-angular detection and thus all angles in one
single measurement [4]. One should note
however that when using angular resolved XPS
(AR-XPS) and the Hf/Si intensities to extract the
130
determination will be essential to understand
their electrical properties. In particular in the
early phase of film growth ("incubation"
period?), a very significant difference can be
observed between the length scale measurements
and the thickness calculated using bulk density
from the number of atoms deposited (Fig. 4a).
For thicker films the two curves show the same
slope and differ by about one nm. Note as
discussed above, that the XPS-results agree with
the RBS data since they are calculated with the
standard attenuation lengths which implicitly
implies the assumption of the normal bulk
density. However when using for each thickness,
the
density
as
derived
from
XRR
(notwithstanding its limitations for very thin
layers), the agreement is much better (Fig. 4b).
The convergence in the early phase of the film
deposition and the disappearance of the 1 nm
offset arises from the unusually low density
values which are measured in that region by
XRR (< 25 % of the bulk value). One must admit
that such a low density is hard to interpret in
terms of a different crystal structure and most
probably must be viewed as indicative for a nonclosed film containing holes or voids
with RBS and the thickness of TEM, quite a
remarkable agreement is obtained (Fig. 3). The
latter holds not only for films with a density
close to the normal bulk values but equally well
for very thin films where the density appears to
be only 25 % of the expected value.
12
CO
I
bulk
phases
>-
(75
LLJ
Q
m
DC
UJ
6
9
12
XRR DENSITY (g/cm3)
Figure 3. Density of HfO2 films (ALD and MOCVD)
as determined from RBS/TEM or XRR. The bulk
values for some phases (C=cubic, M=monoclinic) of
HfO2 are shown for comparison [8].
Compared to the normal bulk density values
(monoclinic, tetragonal or cubic phase of HfO2)
it is clear that high k films with a very different
density can be produced and their density
61-
TEM min/max
SEwithlOOc
RBS He
- TEM min/max
-$~SEwith100c
-A-RBS He with XRR
-0-ARXPS with XRR
-A-XRR
CO
CO
LU
o
I
ALCVD no anneal
25
50
75
100
50
DEPOSITION CYCLES
75
100
DEPOSITION CYCLES
Figure 4. Film thickness versus deposition cycle
a) TEM, SE using optical data derived for the 100 cycles sample, RBS with He beam, single angle and angular resolved
XPS, and XRR fitting both the thickness and the density. RBS/XPS data are converted using bulk density values
b) after recalculation of the RBS and ARXPS data with the XRR-densities for each individual sample [8]
131
The existence of these holes can also be inferred
from the epitaxial alignment of the poly-grains
through the holes when depositing poly-silicon
in-situ on top of a very thin (1.5 nm) high k film
(Fig-6k
Growth mode
Faced with limitations of the growth techniques
it is essential, at least in the early phases of
materials development, to be able to assess in
more detail the detailed nature of the film
growth. A very effective and quantitative way to
look at the film growth in the early stages is to
use Low energy ion scattering (LEIS). In
essence LEIS provides a way to probe the
evolution of the surface coverage of the Sisurface by the growing metal oxide. Hence the
(dis)appearance of the metal(Si)-atoms as a
function of material deposited will be
characteristic for the nature of the film growth
(planar, islanding,..). When thus combining the
film thickness results (RBS or XPS) with the
surface coverage from LEIS (Fig. 5) for the case
of Zr-oxide grown on a HF-last Si-surface, one
can clearly observe the long persistence of a
(small) Si-signal, indicative of Si-areas
remaining uncovered with the metal oxide. The
fact that some Si-atoms are not covered by ZrO2
up to a film thickness in excess of 4 nm, implies
that the some holes in the film remain. Obviously
this will have an impact on the electrical quality
of the film. Such a delayed closure has been
observed for A12O2, HfO2, ZrO2 when deposited
by ALD and is related to the difficulties in
adsorbing on the H-terminated Si-surface versus
the easier adsorption on the growing metal oxide
island [7]. A detailed modeling of this process
has been presented elsewhere [10]. The extent of
this island growth is a function of the surface
preparation and becomes drastically reduced
when growing on "wet" chemical oxides [11].
Figure 6. TEM-cross section of polysilicon deposited
in-situ on thin high k film. The holes in the film allow
the poly to grow epitaxially aligned with the substrate.
As LEIS is rather slow and time consuming an
alternative approach based on TOFSIMS has
been developed for looking at film closure [12].
The technique is based on the fact that
TOFSIMS is also very surface sensitive (the
escape depth of a secondary ion being around 12 monolayers) and hence the Si-signal should
decay very rapidly when becoming covered with
a high k film. For an ideal layer by layer growth,
the Si-intensity would decay by 0.2-0.4 nm/dec.
If the film would grow in a less planar mode, a
slower Si-decay will arise. Hence the steepness
of the Si-intensity decay can be used as a
measure for the efficiency of film closure. An
example of the Si-decay for the deposition of Zroxide on a HF-last surface or an RTO-oxide is
shown in Fig. 7. In both cases the decay of the
Si-intensity is far less than the theoretical 0.2-0.4
nm/dec. In particular the very slow decay for the
HF-last case indicates a very slow closure of the
film, as can deduced from the LEIS-results as
well.
•*~Si for Zr/HF
-*-Si for Zr/RTO
^-Si (LEIS) for Zr/HF
1
2
3
4
5
2
Thickness (nm)
4
XPS-layer thickness (nm)
Figure 5. Si(Zr) surface concentration as determined
with LEIS for ALD-growth of ZrO2 on HF-last Si
surface. Thickness is deduced from XPS
measurements assuming bulk density.
Figure 7. Decay of Si-intensity (-surface
concentration) from TOFSIMS versus high k film
thickness for deposition on RTO and HF-last surface.
Si-surface concentration based on LEIS for the HFlast sample is shown as well.
132
the presence of the metal oxide. Indeed whereas
one would expect a linear decrease of the Siintensity (and a linear increase of the metalintensity) with increasing metal-oxide coverage
(as seen in LEIS), in reality these correlations are
far less linear (Fig. 8). Needless to say that the
non-linearity of these curves will affect the
steepness of the Si-decay in addition to the effect
of a decreasing metal coverage. Hence the decay
curves can not be interpreted as a direct
measurement of surface coverage but should
solely used in a qualitative manner. Certainly
within one materials system, these non-linearities
are fairly constant and a steeper decay of the Siintensity will always be indicative for a faster
closure of the film.
4-AI/rtonoair
A Al/rtoair
«-Hf-HFLast
D HfchoxO.Snm
A HflnmRTO
O--Hf 0.5 nm RTNO
O Hf 0.5 nm RTNO+ H20
• Zr/RTO
-•-Zr/RTO/noair
Figure 8. Si intensity versus metal-intensity as
measured with TOFSIMS for various coverage's of
different metal oxides. Films have been deposited on
different substrates and exposed to air for a short/long
time, prior to the measurement. The straight line
represents the ideal one-to-one relation.
As the evolution from islands towards a closed
film represents the transition from a
heterogeneous surface towards a homogeneous
surface, the growth rate per cycle is continuously
changing as well. Assigning the origin of the
non-linearities in the growth curves to the same
mechanism (delayed closure, heterogeneous
surface) implies that with optimized surface
preparation, whereby films will close more
rapidly, also the growth curves will be more
linear. The latter is evident when comparing the
trend in the RBS growth curves with the
TOFSIMS decays for Hf-oxide films grown on
various surfaces (Fig. 9) [10].
Although in Fig. 7 almost a one-to-one
correspondence between the LEIS and the
TOFSIMS results can be observed, this is not
universally true. When using the TOFSIMS
method to study the growth behavior on different
starting surfaces (Si-OH, Si-H terminated, NH3
annealed surface) as well as for different metal
oxides, one needs to realize that the TOFSIMS
data are only semi-quantitative and possibly
influenced by changes in ionization yield as a
result of the increasing surface oxidation (in
particular in the case of an HF-last surface) or
1.6E+16 •4-HFIast
1.4E+16 •OChem oxide 0.9 nm
i.o -|
1.4
^1nm RTO
1.2E+16 -3-0.5 nm RTNO
1E+16 -«.5nmRTNO + H20
fi
1.2
I
I
\
H>Sichox0.9nm
I
-*-Si 1nm RTO
11
8E+15
-*-SiHfLast
w
\
•e-Si 0.5 nm RTNO
"*-Si 0.5 nm RTNO + H20
^
__J
I
6E+15
4E+15
2E+15
o
5E+15
1E+16
1.5E+16
2E+16
Hf surface coverage (mol/cm2)
Figure 9: Growth curves (RBS) (left) and Si-decay (TOFSIMS) (right) for HfO2-films deposited on various surfaces.
Faster closure (steeper Si-decay) corresponds with a more linear growth curve.
temperatures, more complex systems based on
mixed oxides (aluminates, silicates) have been
proposed, as they appear to have a higher
thermal stability. In ALD these mixed oxides are
grown with alternating cycles of the individual
Composition profiles
Faced with the formation of nano-crystallites in
the high k films, even at moderate anneal
133
Hf, some inconsistencies are present as well. For
instance, whereas the Al signal scales with the
deposition ratio, there is very little influence on
the Hf-signal despite the change in Hfconcentration. Obviously the non-linear response
of the SIMS signals is linked to (unknown)
changes in ionization probability with Hf/Alconcentration. These are probably smaller for Al
than for Hf as the ionization potential (one of the
determining factors for ionization) of Al (5.99
components. As the growth rate per cycle is less
than one monolayer, one expects to obtain a
uniform composition throughout the film.
However it has been observed when growing one
kind of high k material on top of another one,
that the initial growth rate is strongly influenced,
hence some interaction of the individual metals
can not be excluded. An indication of such an
effect is found when growing AlHf-oxides with
different thickness and determining their
composition with RBS or XPS. As these films
are very thin (1-2 nm) neither of these techniques
has sufficient depth resolution to determine the
real depth profile and only an average
composition can be determined. The results
shown in Fig. 10 indicate that for very thin films
the Hf composition exceeds the intended
composition significantly suggesting a strong
initial preference for Hf deposition versus Al.
The latter appears to be a general phenomenon
and dependent on the Hf/Al-ratio (higher initial
deposition of Hf with increased Hf/Al ratio)
104
10
3
C ____ — ™
o 3.
2
o"
0. 9
E 2
8 1.5
15
_
-
-™
_______ .
-*-2:1
1
x ^————————*
<>D——————————n
^——————————*
=
Figure 11. TOFSIMS depth profile for HfAl-oxide
films (cycle ratio 1:2 and 2:1).
-0-1:1
^05
*=°'5
1
Time
-*-1:2
^F
eV) is smaller than the one of Hf ( 7eV) making
the ionization less sensitive to the surface
chemistry.
In order to resolve this problem, use should be
made of a more quantitative technique with very
high depth resolution. An excellent tool for this
application is high-resolution elastic recoil
detection (ERDA) [13]. The latter is the
complement of RBS in the sense that now a high
energy (40-170 MeV) heavy element (I-Au) is
used to elastically scatter light elements out of
the target, which are then analyzed according to
their mass and energy. Using very advanced
magnetic sector analysers, sub-nm depth
resolution can be obtained (at least in the nearsurface regions) while maintaining the
quantification properties of nuclear methods. The
results shown in Fig. 12 relate to a 7 nm HfAlfilm deposited on a nominally 1 nm interfacial
oxide. The gradient of the oxygen profile at the
surface is a good measure for the depth
resolution of the technique. In this experiment it
is less than 0.5 nm/dec. The presence of the
interfacial oxide can quite easily be detected
when comparing the depth of the oxygen profile
versus the depth of the Al and the Hf-profiles.
The difference (~1 nm) determines the thickness
of the interfacial oxide after high k film growth.
Consistent with the results in Fig. 10 and 11, the
0- ———————————————————————
0
2
4
200 400 600 800 1
6
thickness (nm)
Figure 10. Average Hf/Al composition ratio versus
film thickness for three "bulk" Hf7Al-cycle ratios (1:2,
1:1, 2:1). As the growth rate per cycle is different for
Hf and Al, the cycle ratios do not reflect the
composition.
The results in Fig. 10 indicate a transient in the
deposition efficiency of Hf versus Al with a
higher Hf- or a reduced Al-deposition in the
early stages of the film growth. Inevitably this
will then lead to a non-homogenous film
composition. In order to be able to probe such a
compositional variation, one needs a technique
with sub-nm depth resolution. One possible
approach is to make a TOFSIMS depth profile
with a very low energy (< 500 eV) sputter beam.
The results in Fig. 11 indicate a strong nonuniformity with a Hf enrichment at the SiOiinterface and a sloped profile for the Hf and the
Al. Although the TOFSIMS profiles show the
expected behavior i.e interfacial enrichment of
134
ERDA-data show a sloped profile for the Hf as
well as for the Al, highlighting the non-linear
growth rates as a function of film thickness.
Quite surprising, but very clear, is the fact that
the first monolayer is completely free of Hf
suggesting that during the growth of the mixed
oxide a phase separation occurs leading to some
Al-floating on the growing surface. The different
slopes (positive versus negative) of the two
elements also point into the direction of a surface
migration of Al and Hf-indiffusion causing an
(extra) enhanced Hf-concentration at the
interface. It is also interesting to look at some
other contaminants such as H and C. The Cprofile solely peaks at the surface and is thus
representative of the contamination adsorbed on
the sample surface during exposure to the
ambient. The H-profile has a double peak
structure : one large peak concentrated at the
surface which is also indicative of the surface
contamination and a smaller one (-peak
concentration (~1 at %) located at the metaloxide- silicon oxide interface. Clearly the latter
is associated with remnants from the water
pulses of the ALD-process, probably reflecting a
poorer removal of OH-groups from the silicon
oxide surface by the N2-pulse than from the
subsequent metal oxide surface.
influenced by the presence of the high k film, an
effect which is reduced with increasing
interfacial oxide thickness. Hence some
interfacial oxide might be required after all,
necessitating appropriate metrology as well. For
this purpose one can rely on XPS, which gives
directly the amount of oxidized Si.
For instance, whereas recipes for the formation
of very thin (< .4 nm) oxide layers as starting
surface have been developed, the interfacial
oxide after the full processing might be much
larger. One of the determining factors is, for
instance, the post-deposition anneal which,
depending on the ambient (inert, oxygen,
nitrogen,..), can cause a significant interfacial
oxidation. The partial oxygen pressure during
this aneal is crucial (Fig. 13).
^ 2.5
^
2
1O
-RTON-800
5
-RTON650
I °•§
o
1.E-04
1.E-02
1.E+00
1.E+02
oxygen pressure (Torr)
Figure 13. Interfacial oxide thickness, as measured
with XPS, for HfO2-fllms grown on 0.4 nm
oxy(nitri)de. Parameter is the oxygen partial pressure
during the anneal.
-2
0
2
4
6
8
10
12
At low pressure (< 10~2 Torr) the interfacial
oxide is limited by the oxygen supply while at
high pressures (> 1 Torr), the interface reaction
becomes the limiting step. The oxygen itself is
highly mobile through the high k film such that
the interfacial oxide is independent of the film
thickness and solely determined by the
temperature dependence of the interface reaction
(Fig. 14). Noteworthy also is that the final oxide
thickness is much larger (~ 2 nm) than what can
be observed on Si annealed without the high k
film present. Obviously the presence of the high
k film has a catalytic effect on the Si-oxidation.
Of course the XPS approach breaks down when
the high k gate stack intentionally contains
silicates. In that case TEM imaging becomes the
only resource for an accurate measurement.
Another word of caution is required with respect
to the interpretation of the XPS-interface
thickness data. It has, for instance, been observed
with TOFSIMS depth profiles (Fig. 15) that
when annealing mixed oxides, an element
14
depth (nm)
Figure 12. ERDA-profile of a HfAl-oxide on a 1 nm
SiO2/Si substrate.
INTERFACIAL OXIDE
One of the crucial points in obtaining a high
quality high k film meeting the (very) low EOTrequirements (< 1 nm) is the control of the
interfacial oxide. Ideally no oxide layer should
be present as this impacts on the available
thickness budget of the high k film [6]. However
the presence of an oxide has a positive impact on
the film growth (more linear growth curves,
more rapid closure). Last, but not least, it has
been established that the mobility is negatively
135
characterization methods at the macroscopic
device level. These tests, however, provide only
spatially averaged information on the electrical
properties of the material and do not address the
microscopic failure mechanisms. One very
interesting approach to address this issue is to
use Conductive Atomic Force Microscopy (CAFM), as already demonstrated for SiO2 [14].
Conductive Atomic Force Microscopy (C-AFM)
is a measurement technique based on an AFM
equipped with a conductive tip, whereby
simultaneously topographical and electrical
(tunneling current) information is collected with
a typical lateral resolution of about lOnm.
Moreover,
local
current-voltage
(I-V)
characteristics can also be measured in a range
from 100 fA to 100 pA. As the current tunneling
from the tip through the oxide is strongly
dependent on the film properties, leakage paths
and trapped charges, C-AFM is an excellent tool
to study on a local scale the electrical quality of
the film and the breakdown mechanisms.
4L.9
2
?
-*-8
8
(75
•o
—, ^
-•-7 OOC, N2O
-*-no anneal
y
^^^M
\T
0.5
A
V
0
1
1
20
40
60
film thickness (cycles)
Figure 14 : Interfacial oxide versus film thickness of
mixed oxide before and after anneal.
migration occurs during the anneal whereby
some Si diffuses towards the surface. Needless to
say that in that case the apparent interfacial oxide
will be overestimated as the photoelectrons from
the surface oxide are far less attenuated than the
ones from the deeper lying interface. Note that
concurrent with the Si-diffusion migration of Al
and Hf can be observed.
I
5o
200
400
600
800
1000
1200
Sputter Time (arb. unit)
Figure 15.TOFSIMS depth profile of HfAl-oxide
before and after anneal. Arrows indicate the element
migration with increasing anneal temperature.
Figure 16. Current maps (-6 V bias) of HfAl- gate
stack for different annealing temperature :. a) not
annealed sample, b) TA=800C, c) TA=900C, d)
TA=1000C. The area of the scanned region is 2x2jim2.
The current ranges from Op A (black) to 20pA (white),
e) 1x1 urn2 zoom of a), showing weak (more
conductive) spots, when the current scale is decreased
from OpA (black) to IpA (white), f) a 0.3x0.3|im2
current image of the sample exposed to TA>900C. The
current ranges from OpA (black) to lOpA (white) [15].
LOCAL ELECTRICAL
PROPERTIES
Most of the knowledge about the electrical
behavior of high-k materials has been gained
from measurements performed on macroscopic
MOS capacitors or transistors using standard
136
understand the details of film growth and the
film composition. The measurement of the local
electrical properties with C-AFM enables one to
link local electrical deficiencies (weak spots)
with structural defects or nanocrystallites.
In the example shown in Fig. 16, HfAl-oxide
films have been measured as a function of their
post-deposition anneal temperature [15].
Basically the current maps show a low
homogeneous leakage current distribution for the
as-deposited samples with a gradually increasing
leakage current for higher anneal temperatures.
At the same time one observes the appearance of
local regions of even higher conductivity.
Obviously the latter represent "weak" spots (not
yet breakdown spots!) in the film which may
evolve into complete film failure. The size of
these defects is close to the tip resolution and on
the order of 5-20 nm. Their density increases
with anneal temperature and can become as large
1010 defects/cm2. Although their nature is not
complete understood yet, it is tempting to
correlate these spots with the formation of nanocrystallites in the high k film as observed in
TEM-images (Fig. 17) and increased conduction
through these grains or their grain boundaries.
REFERENCES
1. D.A. Cole, J.R. Shallenberger, S.W. Novak,
R.L. Moore, M.J. Edgell, S.P. Smith, CJ.
Hitzman, J.F. Kirchhoff, E. Prinicpe, W.
Nieveen, F.K. Huang, S. Biswas, R.J. Bleiler and
K Jones, J. Vac. Sci. Technol. B 18, 440-444
(2000).
2. W.W.Chism, A.CDiebold and J.M.Price,
these proceedings
3. P.Eoher, P.Evrard, C.Defrenoux, J.C.Fouere,
E.Bellandi and H.Bender, these proceedings
4. R.Champanaria, P.Mack, R.White and
J.Wolstenholme, these proceedings
5. P.J. Cumpson, Surf. Interface Anal. 29, 403406 (2000).
6. J.Wolstenholme, presented at the Fall
meeting of the American Vacuum Society,
Denver, 2002.
7. W.Vandervorst, B. Brijs, H. Bender, O.T.
Conard, J.Petry, O.Richard, S. Van Elshocht, A.
Delabie, M. Caymax and S. De Gendt in Novel
Materials and Processes for Advanced CMOS,
edited by M. Gardner, S. De Gendt, JP Maria and
S. Stemmer, (Mater. Res. Soc. Proc, Pittsburgh,
USA), 745, 23-34 (2003)
8. H. Bender, Th. Conard, O. Richard, B. Brijs,
J. Petry, W. Vandervorst, Chr. Defranoux, P.
Boher, N. Rochat, C. Wyon, P. Mack, J.
Wolstenholme, R. Vitchev, L. Houssiau, J-J.
Pireaux, A. Bergmaier, G. Dollinger, The
Electrochemical Society,
Proc . ALTECH,
Paris-2003.
9. Matty Caymax, H. Bender, B. Brijs, T.
Conard, S. De Gendt, A. Delabie, M. Heyns, B.
Onsia, L.Ragnarsson,
O. Richard, W.
Vandervorst, S. Van Elshocht, C. Zhao, J.W.
Maes, L. Date, D. Pique, E. Young, W. Tsai, Y.
Shimamoto, Proc. Mater.Res.Soc, Spring
Meeting-2003, San Francisco
10. M.A.Alarn,
M.L.Green,
M.Y.Ho,
W.Vandervorst, B.Brijs, T.Conard, P.LRaisanen,
presented at Mater. Res. Symp, symposium N
(Novel Materials and Processes for Advanced
CMOS), Fall-2003, Boston, USA
11. M. L. Green,et al. J.App. Phys 92, 71687174(2002)
Figure 17. Cross-sectional TEM image of a sample
annealed at 900C. Crystalline grains (ellipse) can be
observed.
CONCLUSIONS
The study of the growth of high k dielectrics
shows that the formation of very thin high
quality films does not occur easily with ALD. At
present in ALD, the growth process can be
highly non-linear and film closure is strongly
dependent on the surface preparation. This 3Dgrowth requires the deposition of rather thick
films (3-4 nm) before completely closed films
are obtained. Therefore film thickness metrology
becomes an important issue whereby one needs
to distinguish between the physical thickness (as
measured with SE/TEM) and the areal density
(as obtained with XPS/RBS). The comparison
between both leads to the evaluation of the film
density, which for very thin films deviates
drastically from the bulk values. High resolution
techniques like ERDA and TEM are essential to
137
12. T.Conard, W.Vandervorst, J.Petry, C.Zhao,
W.Besling, H.Nohira and O.Richard,
Appl.
Surf. Sci. 203-204,400 (2003)
13. G.Dollinger.
C.M.Frey,
A.Bergmaier,
T.Faestermann, Nucl. Instrum.Meth.B 136-138
603 (1998)
14. M. Porti, M. Nafria, X. Aymerich, A.
Olbrich, and B. Ebersberger, Appl. Phys. Lett. 78
4181-4183(2001).
15. X. Blasco, J. Petry, M. Nafria, X.
Aymerich, W. Vandervorst, to be published in
proc INFOS-2003, Barcelona, Spain.
138