Phase formation of B2-RuAl during annealing of Ru/Al multilayers

Intermetallics 18 (2010) 1507e1516
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Intermetallics
journal homepage: www.elsevier.com/locate/intermet
Phase formation of B2-RuAl during annealing of Ru/Al multilayers
N. Zotov a, *, K. Woll b, F. Mücklich b
a
b
Institute of Materials, Department of Mechanical Engineering, Ruhr University Bochum, D-44780 Bochum, Germany
Chair of Functional Materials, Department of Materials Science, Saarland University, D-66041 Saarbrücken, Germany
a r t i c l e i n f o
a b s t r a c t
Article history:
Received 10 November 2009
Received in revised form
31 March 2010
Accepted 1 April 2010
Available online 28 April 2010
The formation of B2-RuAl from Ru/Al multilayers (MLs) with an average MLs composition of Ru47Al53 and
modulation periods L up to 22.4 nm was studied by in-situ X-ray diffraction (XRD), differential scanning
calorimetry, scanning electron microscopy and transmission electron microscopy. The as-deposited MLs
with L < 4.5 nm grow epitaxially with relatively small roughness of the atomic layers. At higher L values,
the epitaxy is lost and polycrystalline MLs with strongly distorted atomic layers develop during deposition. In-situ high-temperature XRD demonstrated that L influences the phase evolution and kinetics
during annealing. At annealing temperatures TA < 673 K Al diffuses into the Ru layers leading to the
formation first of Ru(Al) solid solution. At TA > 823 K the ordered B2-RuAl phase is formed via a diffusioncontrolled nucleation. The RuAl grain growth kinetics accelerates with increasing L.
Ó 2010 Elsevier Ltd. All rights reserved.
Keywords:
A. Aluminides, miscellaneous
B. Diffusion
C. Thin films
D. Phase interfaces
F. Diffraction
1. Introduction
The exceptional high-temperature strength as well as the high
creep and oxidation resistance of B2 type intermetallic aluminides
makes this group of intermetallics interesting for high-temperature applications [1e4]. Possible applications are oxidation
protection scales in thermal barrier systems [5] or protective thin
films which avoid material interactions at high temperatures, e.g.
in moulding dies [6,7]. However, a limiting factor is the brittleness
of intermetallic aluminides at room temperature (RT) as they show
a brittleeductile temperature TBD, above 773 K [8]. A comparative
study of many intermetallics with melting points above 1973 K
reveals four binary systems that demonstrate exceptional room
temperature ductility. One example is RuAl [9]. It is characterized
by high ductility at RT, as compared with aluminides like NiAl, CoAl
or PdAl.
Due to its extraordinarily high melting point of 2333 K, an
excellent thermodynamic stability is guaranteed up to high
temperatures. Good high-temperature strength is known and high
creep resistance can be expected. Recent experiments indicated
that B2-RuAl also shows excellent oxidation properties due to the
formation of a thin protective Al2O3 scale [2,10]. In addition, Tryon
et al. [11] reported that the thermal expansion coefficient of B2RuAl is nearly equal to that of Al2O3, whereas the thermal
* Corresponding author. Tel.: þ49 234 3229109; fax: þ49 234 3212435.
E-mail address: [email protected] (N. Zotov).
0966-9795/$ e see front matter Ó 2010 Elsevier Ltd. All rights reserved.
doi:10.1016/j.intermet.2010.04.001
expansion coefficient of B2-NiAl exceeds that of alumina about one
order of magnitude. Hence, thermal-induced residual stresses at
the interface between alumina and B2-RuAl are expected to be
lower and crack initiation due to thermal load should be significantly reduced.
For all of the potential applications mentioned above RuAl
coatings or thin films are generally needed. Thermal annealing of
MLs is often used as a fast reaction pathway for the preparation of
new thin film structures [12] due to the high density of defects and
diffusing species at the interfaces, which can lead to changes in the
effective heats of formation [13]. Thus, typical reaction temperatures between the elemental layers in MLs are significant lower
than 1000 K.
The present study is focused on the formation of RuAl intermetallic thin films by thermal annealing of Ru/Al multilayers (MLs).
Up to now, reactions in Ru/Al MLs were studied in equiatomic MLs
with modulation period L ¼ 88 nm [14] (L ¼ tRu þ tAl, where tRu and
tAl are the thicknesses of the corresponding atomic layers).
Annealing of these MLs at low temperatures results in the formation of RuAl6. Transformation to the B2-RuAl phase occurs only for
annealing temperatures TA > 873 K. Studies on the reaction
behavior of similar MLs, e.g. Ni/Al [15] and Ti/Al [16], reveal,
however, a strong influence of L on the reaction behavior. These
studies demonstrated that MLs with small L (LNi/Al ¼ 5 nm, LTi/
Al ¼ 4 nm) transform directly to the equilibrium phases during
annealing. Hence, concerning potential applications where monolithic RuAl thin films are needed, small L MLs should be used.
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N. Zotov et al. / Intermetallics 18 (2010) 1507e1516
Consequently, the main goals of the present paper are the
detailed investigation of thermal stability and reaction behavior in
Ru/Al MLs with small periods in the range 2.2 nm < L < 22.4 nm,
the formation of B2-RuAl from such MLs as well as better understanding of diffusion and phase formation in metallic multilayers in
general.
lamellae with typical thicknesses 100 nm were prepared from
each sample using focused ion beam (FIB) technique (Dual Beam
Helios Nanolab, FEI). Pt cap layers were deposited before the FIB
preparation to avoid damage of the MLs. For image analysis of the
SEM and TEM micrographs, both the ImageJ Ver. 1.41 [17] and the
a4i Analysis (Soft Imaging System GmbH) software were used.
2. Experimental details
2.2.4. Differential scanning calorimetry
Differential scanning calorimetry (DSC) measurements were
performed on selected samples using Perkin Elmer DSC 7 thermal
analyzer. The thermal analyzer was calibrated against the known
melting temperatures and enthalpies of Zn, In and Pd standards.
The DSC curves were measured at a heating rate of 20 K min1
under Ar heat flow. Two subsequent runs for each sample were
made to better estimate the baselines. For the DSC measurements,
MLs were deposited on mica substrates and peeled off. The peeled,
free-standing films (w5 mg weight) were loaded into Al pans.
Ru/Al MLs with nominal compositions [Ru x/Al 1.244x]n, where
x ¼ 1, 2, 3, 4, 5 and 10 nm, were deposited on (100) Si wafers by DC
magnetron sputtering using elemental high-purity Ru and Al
targets in a von Ardenne PPS-A 200 system. No buffer-layer was
deposited between the substrate and the MLs. The base pressure
was 0.1 mPa. The Ar gas sputtering pressure was 0.3 Pa. To
construct the MLs, the substrate was alternately rotated under the
corresponding target. The Ru sputtering rate was 0.2 nm/s, the Al
sputtering rate was 0.24 nm/s. The number of bilayers n in the MLs
(n ¼ 80, 40, 27, 20, 16 and 10, respectively) was selected so that the
nominal total thicknesses of the MLs to be between 177 and
230 nm. The average chemical composition of all MLs e Ru47Al53,
was determined by energy dispersive X-ray spectroscopy using
Dual Beam Workstation (Strata DB 235) operating at 10 kV with
120 s accumulation time.
2.2. Structural characterization
2.2.1. X-ray diffraction
All as-prepared MLs were first measured at RT by X-ray
diffraction (XRD). A Bruker Discovery D8 diffractometer and Cu Ka
radiation (40 kV, 40 mA) were used. It was equipped with a twodimensional (2D) ‘Histar’ detector able of measuring a large 2q
range (here about 30 ). The measuring time was 1200e1800 s/
frame. To study the phase evolution in the MLs during annealing,
in-situ high-temperature (HT) XRD measurements were performed
on selected samples. For this purpose, a special heating stage from
MRI GmbH mounted on the D8 diffractometer was used. The
samples were mounted on a BN-coated graphite heater and pressed
against the surface of the heater with special metal frames to
achieve better mechanical stability and better thermal contact. The
water-cooled heating stage was covered with a Be hemisphere
(100 mm diameter, 0.5 mm thickness) and evacuated. The averaged
pressure in the chamber was 9 Pa. The temperature was measured
using a type-K thermocouple inserted in a hole in the heater, close
to the hot zone, and controlled with accuracy of about 3 K.
Patterns were measured at RT, 473, 598, 828 and 873 K, selected on
the basis of differential scanning calorimetry (DSC) measurements.
The measuring time for each pattern was 10 min. The XRD patterns
were corrected afterwards for background scattering from the Be
dome and Fourier smoothed, where necessary.
2.2.2. Scanning electron microscopy
The in-plane microstructure of the as-deposited MLs was
investigated using a Leo 1530 VP scanning electron microscope
(SEM) with field-emission gun. SEM micrographs were taken at
5 keV using an Inlens detector. Magnifications of 20,000, 40,000
and 80,000 were chosen.
2.2.3. Transmission electron microscopy
Transmission electron microscopy (TEM) images of the [Ru
1 nm/Al 1.24 nm]80, [Ru 4 nm/Al 5 nm]20, [Ru 5 nm/Al 6.22 nm]16
and the [Ru 10 nm/Al 12.44 nm]10 MLs were taken using Jeol 2010
TEM, operating at 200 kV, using different magnifications. TEM
3. Results and discussion
The phase formation during annealing of MLs depends sensitively on their microstructure and composition. Therefore, the
microstructure of the investigated Ru/Al MLs is characterized by
XRD, SEM and TEM in Sections 3.1e3.3.
3.1. X-ray diffraction at room temperature
Reconnaissance XRD measurements showed that under
symmetric reflexion conditions the MLs scatter strongly mainly in
the 2Q range of 30e55 . The corresponding XRD patterns (corrected for substrate scattering) are given in Fig. 1. The investigated
MLs show rather complex structural behavior with increasing L.
The XRD patterns of the MLs with L 4.5 nm show only one very
broad and asymmetric peak at about 42e43 2Q with some small
intensity oscillations. In contrast, the XRD patterns of MLs with
L 4.5 nm exhibit 3 broad diffraction peaks at about 39, 42 and
44 2Q.
The diffraction patterns from strongly-textured crystalline MLs
(also called superlattices) with scattering vector Q perpendicular to
the multilayer’s surface (Q ¼ 4p sin(Q)/l), where l is the X-ray
wavelength used, are characterized by Bragg peaks of the average
structure, centered at Q00l ¼ (2p/do)l, l ¼ 0, 1, 2,. (where do is the
average interplanar distance along the growth direction) and
satellites around them, arising from the compositional modulation
of the MLs. For MLs without variations of the interplanar spacings
10000
8000
Intensity (arb. units)
2.1. Sample preparation
6000
Λ = 22.44 nm
4000
2000
0
32
11.22
9.00
6.73
4.48
2.24
34
36
38
40
42
44
46
48
50
52
2Θ (degrees)
Fig. 1. X-ray diffraction patterns of the as-deposited MLs in the range 32e52 2Q.
N. Zotov et al. / Intermetallics 18 (2010) 1507e1516
Table 1
Simulated structural parameters of the MLs with periods L < 4.5 nm.
Lnom (nm)
tAl (nm)
dAl (nm)
tRu (nm)
dRu (nm)
Lexp (nm)
s(L)a (nm)
2.24
4.48
1.1
2.6
0.23
0.235
1.05
1.9
0.212
0.212
2.15
4.5
0.10
0.13
a
Random variations in the modulation period L.
and/or the superlattice period L, the positions of the satellites are at
Qm ¼ Q00l (2p/L)m, m ¼ 1, 2,. [18].
Initially, simulations of the diffraction profiles in the kinematic
approximation were done for all Ru/Al MLs using the nominal layer
thicknesses and no L-disorder. It was further assumed that facecentered cubic (fcc) Al and hexagonal close packed (hcp) Ru layers
are formed during the deposition, that the Al layers grow along the
(111) fcc direction, while the Ru layers grow along the (0001) hcp
direction. Both crystallographic planes are the planes with largest
in-plane atomic densities and are most commonly formed during
physical vapor deposition. For example, Wen et al. [19] observed
a growth of the Ru layers along the (0001) direction in Ni/Ru MLs.
Correspondingly, the d-spacings of the Al and Ru layers in these
simulations were taken equal to dAl
(111) ¼ 0.234 nm [20] and
dRu
(0002) ¼ 0.214 nm [21], respectively. These simulations showed
that the XRD patterns at RT cannot be modeled as ideal superlattices and the peaks in the XRD patterns of MLs with L > 4.5 nm
do not correspond to ML satellites.
A reasonable agreement between simulations and experiment
for the MLs with L 4.5 nm was achieved by adjusting the thicknesses and the d-spacings of the Ru and Al layers as well as introducing Gaussian L-disorder. The results of these simulations are
given in Table 1 and shown in Fig. 2 for L ¼ 4.48 nm. No distinct
satellite peaks are observed due to the presence of L-disorder.
Nevertheless, the ability to simulate the diffraction patterns of the
MLs with L < 4.5 nm, indicates that Ru and Al grow by coherent
epitaxy on each other, similar to the Ru/Ni MLs with L < 20 nm [19].
The calculated superlattice periods are close to the nominal ones.
That is why the samples will be further denoted with the corresponding nominal L. The d-spacings of the Ru layers are slightly
smaller than the bulk dRu
(0002) value.
The decrease of dRu
(0002) could be related to the presence of inplane tensile stresses in the Ru layers. The presence of L-disorder
suggests some chemical (due to diffusional processes) and/or
morphological roughness of the Ru/Al and Al/Ru interfaces.
The XRD patterns of the samples with L > 4.5 nm were
decomposed using 3 pseudo-Voigt functions [22], taking into
Fig. 2. X-ray diffraction pattern of the as-deposited ML with L ¼ 4.48 nm (points) and
simulation (full line).
1509
account the presence of Ka1 and Ka2 components for each peak. A
typical fit is given in Fig. 3a for L ¼ 9.0 nm. The d-spacings of the
three peaks for all samples are in the ranges 0.230e0.234 nm,
0.212e0.216 nm and 0.205e0.209 nm, respectively. These peaks
can be assigned to the Ru(10e10)/Al(111), Ru (0002) and Ru
(10e11) peaks, respectively [20,21]. The goodness-of-fit factors
were between 2% and 5%. HT XRD measurements (Section 3.4.2)
showed that the peak at about 38.5 2Q arises mainly from Al(111).
All these results suggest that the MLs with L > 4.5 nm are polycrystalline. This is further confirmed by measurements at higher 2Q
angles (Fig. 3b) showing the (10e12), (11e20), (31e41) Ru as well as
weak (220) and (311) Al peaks.
Texture often influences the grain growth during annealing (see
for example the work of Semiatin et al. [23] on TieAleV intermetallics). That is why, the possible presence of preferred orientation
in the investigated MLs was checked by recording diffraction
patterns of all samples for different angles c between the normal to
the thin film surface and the diffraction plane (for a symmetric
Bragg-Brentano geometry c ¼ 0 ). Change of c for the MLs with
L < 4.5 nm does not lead to significant changes in the intensities
(see Fig. 4a), suggesting relatively broad distribution of the
misorientations of the MLs grains from the thin film normal. This is
also confirmed by inspection of the 2D images from the Histar
detector. Fig. 4b shows XRD patterns at two c angles for the
Fig. 3. (a) X-ray diffraction pattern of the as-deposited ML with L ¼ 9.0 nm (points) in
the range 30e52 2Q, profile fit (thick line), individual peaks (dashed lines) and
difference plot (thin line). (b) X-ray diffraction pattern of the as-deposited ML with
L ¼ 9.0 nm (points) in the range 55e82 2Q and profile fit (thick line).
1510
N. Zotov et al. / Intermetallics 18 (2010) 1507e1516
Table 2
Grain sizes L0002 and microstrains e0002 of the Ru phase in the MLs with L > 4.48 nm.
L (nm)
L0002 (nm)
e002
6.73
9.00
11.22
22.44
6.6 0.7
7.9 0.5
7.8 1.0
10.9 1.5
0.02 0.02
0.016 0.005
0.016 0.005
0.012 0.008
a tendency for a decrease of the microstrains with increasing L.
LRu
0002 is larger than the nominal thickness of the Ru layers for the
MLs with L > 6.73 nm. Presumably, the Ru layers make intergrowths and/or short-cuts interrupting the ideal ML structure.
3.2. Scanning electron microscopy of the as-deposited MLs
The in-plane microstructures of the [Ru 1/Al 1.24]80 and [Ru 2/Al
2.48]40 MLs are similar and consist of closely-packed almost
circular grains (see Fig. 5a). Visually, the intergranular space in
these MLs is small with typical thickness of grain boundaries (GB)
between grains in close contact of about 15 5 nm. The in-plane
microstructures of the [Ru 5/Al 6.22]16 and [Ru 10/Al 12.44]10 MLs
are quite different. Typical chain-like agglomerates, containing 3e6
grains, are formed in these MLs (Fig. 5b). These agglomerates are
separated by triple junctions and intergranular space.
In order to quantify the changes in the in-plane microstructures
with increasing L, the average grain sizes of well-separated
Fig. 4. (a) X-ray diffraction patterns of the as-deposited ML with L ¼ 4.48 nm at
different c angles. (b) X-ray diffraction patterns of the as-deposited ML with
L ¼ 22.44 nm at different c angles.
polycrystalline ML with L ¼ 22.44 nm. There is nearly no intensity
variation of the Al (111)/Ru (100) peak. However, the Ru (0002)
peak decreases whereas the Ru(10e11) peak increases with
increasing c. The same behavior is observed for all MLs with
L > 4.5 nm. It can be concluded that on the one hand the Al (111)
grains have relatively broad distribution of misorientations from
the thin films normal. On the other hand, the Ru layers are strongly
textured along the (0001) direction.
A decrease of the full width at half maximum of the Al (111), Ru
(0002) and Ru (10e11) peaks is observed with increasing L. A grain
size increase, perpendicular to the thin film surface, can therefore
be expected. The average grain size LRu
0002 and the average microstrains eRu
0002 along the Ru (0001) direction were estimated from
single-line Voigt analysis [24] and are given in Table 2. More
specifically, the Al (111), Ru (0002) and Ru (10e11) diffraction peaks
were first fitted by pseudo-Voigt functions [22]. The integral widths
of the Cauchy (bCh) and the Gaussian (bGh) components of the
measured profiles (h) were then calculated using expressions given
in Ref. [24]. The integral widths bCg and bGg of the instrumental
profile (g) were calculated in the same way. A corundum standard
from the United States National Institute of Standards was used for
determination of the instrumental profiles. The integral widths of
the ‘true’ structurally broadened profiles (f) were then calculated as
[24]: bCf ¼ bCh bCg, b2Gf ¼ b2Gh b2Gg. Finally, the grain size LRu
0002 and
Ru
0002
the microstrain eRu
0002 were obtained as: L0002 ¼ l cos(Q0002)/bCf
0002
Ru
and eRu
0002 ¼ ¼bGf ctg(Q0002). L0002 increases, while there is
Fig. 5. SEM micrographs of the as-deposited: (a) [Ru 1 nm/Al 1.24 nm]80 and (b) [Ru
10 nm/Al 12.44 nm]10 MLs.
N. Zotov et al. / Intermetallics 18 (2010) 1507e1516
individual grains and the fraction of intergranular area were estimated using image analysis. Several SEM micrographs with
different magnifications were analyzed for each ML. The average inplane grain size hAki increase with L (Fig. 6a). The fraction of
intergranular space also increases initially with L and then starts to
level-off for L > 11.24 nm (Fig. 6b). The standard deviations of both
hAki and the fraction of intergranular space are relatively big,
despite the large number of grains/voids analyzed, as the corresponding distributions are quite broad and asymmetric. In addition,
the SEM micrographs suggest that the intergranular channels
become not only larger but also deeper with increasing L.
1511
As expected from the XRD results, systematic changes in the
microstructure are observed by TEM with increasing L. The atomic
layers of the sample with L ¼ 2.24 nm are extending over large
number of columnar grains (Fig. 7a), only a few lateral defects are
observed and the columnar grains are separated predominantly by
low-angle GBs. The variation of the modulation periods, estimated
from the TEM images, is small (w0.3 nm) in qualitative agreement
with the XRD results (Table 1). This explains why it was possible to
3.3. Transmission electron microscopy of the as-deposited MLs
Bright field TEM cross-sectional images of the [Ru 1 nm/Al
1.24 nm]80, [Ru 4 nm/Al 5.0 nm]20 and the [Ru 10 nm/Al
12.44 nm]10 MLs are shown in Fig. 7. The TEM images of the [Ru
5 nm/Al 6.22 nm]16 ML are similar to that of the [Ru 4 nm/Al
5 nm]20 ML. The periodic modulation of the MLs and a columnar
microstructure are observed for all samples. Growth of columns,
perpendicular to the substrate, was also observed in Ni/Ru [19], Ag/
Cu [25], Ni80Fe20/Cu [26], NiFe/Ag [27] as well as Co/Cu [28] MLs, for
example. The width of the columns close to the surface increases
from 18 5 nm for the [Ru 1 nm/Al 1.24 nm]80 ML to 50 12 nm for
the MLs with L 9.0 nm, in qualitative agreement with the
behavior of the in-plane grain sizes, determined by SEM (Fig. 6a).
In-Plane Grain Area (nm2)
a
4000
3000
2000
1000
0
0
5
10
15
20
25
Modulation Period Λ (nm)
b
35
Intergranular Area (%)
30
25
20
15
10
5
80 kX
40 kX
0
0
5
10
15
20
25
Modulation Period Λ (nm)
Fig. 6. Variation with L of: (a) the average grain area and (b) the intergranular fraction,
determined by image analysis of SEM micrographs.
Fig. 7. Bright field TEM cross-sectional images of the as-deposited MLs with (a)
L ¼ 2.24, (b) 9.0 and (c) 22.44 nm. Al-rich phase (gray layers), Ru-rich phase (dark
layers).
N. Zotov et al. / Intermetallics 18 (2010) 1507e1516
Fig. 8. Schematic representation of the microstructural characteristics of the MLs with
L > 4.48 nm. Ru/Al intermixed areas are not shown for clarity.
8
6
4
5.0
Λ = 2.24 nm
2.5
ΔZ (nm)
model the XRD patterns of the MLs with L 4.48 nm as coherent
superlattices (Fig. 2). With increasing L, the amount of disorder in
the ML structures increases significantly (Fig. 7b and c). The atomic
layers on the left and the right side of GBs are often shifted vertically. Larger number of discontinuous atomic layers is clearly
observed. They are separated by high-angle GBs, bridges connecting neighboring atomic layers of the same kind and strongly
interdiffused areas. A typical microstructure of MLs with
L > 4.48 nm showing the main defect features is schematically
given in Fig. 8. The average modulation periods determined from
TEM are similar to the nominal ones. But the thicknesses of the
atomic layers of the MLs with L > 4.48 nm start locally to vary
significantly. This is reflected in the standard deviation s(L),
determined from TEM, which increases from 0.3 nm for the [Ru
1 nm/Al 1.24 nm]80 ML to 0.6 nm for the [Ru 4 nm/Al 5.0 nm]20 ML
and to 2.7 nm for the [Ru 10 nm/Al 12.44 nm]10 ML. Altogether,
these microstructural features lead to an increase of the roughness
and a decrease of the lateral correlation length of the atomic layers.
In order to quantify the changes in the microstructure observed
by TEM, the roughness of the Ru atomic layers (sRu) was determined using image analysis of the TEM micrographs. Three
different atomic profiles with lengths between 100 to 400 nm for
each sample were analyzed and the sRu results averaged. A mean
center line for each Ru layer was first defined and then the
roughness calculated as the root-mean-square deviations of the Ru
profile from this line along the growth direction of the MLs. A
typical Ru profile for the ML with L ¼ 2.24 nm is shown as an inset
in Fig. 9. Since the roughness of the Al layers generally replicates
the roughness of the Ru layers, Fig. 9 shows that the roughness of
the MLs increase with increasing L. This is consistent with the
observation that the grooving of the ML surface increases visually
with L.
One possible explanation for the observed rough ML microstructures could be an island-type Ru growth on the top of the Al
layers. An island-type Ru growth can be qualitatively explained in
the following way. Assuming a substrate temperature TS w 353 K,
the homologous temperature Th ¼ TS/Tmelting for Al is 0.38 whereas
for Ru it is 0.14. At such Th, in accordance with the structure zone
models of Movchan et al. [29] and Thornton [30], a Zone T or Zone II
type-of-growth for Al can be assumed. In these zones the arriving
Al atoms are sufficiently mobile due to significant surface diffusion.
The Al atoms can therefore migrate on the surface before being
covered by the next atomic layer. A relatively flat Al surface should
be the result (see the first Al layers next to the substrates in Fig. 7).
However, the Ru-layers’ growth should be in the Zone I mode
[29,30] due to the small Th for Ru. Therefore, the mobility of the Ru
adatoms is expected to be low. This reduced mobility of the Ru
Ru Layer Roughness (nm)
1512
2
0.0
-2.5
-5.0
0
20
40
60
80
100
Latheral Length (nm)
0
0
5
10
15
20
25
Modulation Length Λ (nm)
Fig. 9. Variation with L of the Ru atomic layers roughness, determined from TEM
(details see text).
atoms favors local grain growth [31], perpendicular to the
substrate, and island-type morphology of the Ru layers might be
expected.
3.4. Phase formation at elevated temperatures
Typical DSC curve of the ML with L ¼ 22.44 nm is shown in
Fig. 10. Two broad exothermic peaks are observed at about 573 and
773 K. Correspondingly, in-situ HT XRD measurements were performed at selected TA below the fist DSC peak (RT, 473 K), between
the two DSC peaks (698 K) and above the second DSC peak (823 and
873 K), in order to determine the phases formed.
3.4.1. Multilayers with L < 4.5 nm
Fig. 11 shows the HT XRD patterns of the ML with L ¼ 4.48 nm.
The XRD patterns were Fourier smoothed. Annealing up to
TA ¼ 698 K leads only to minor changes in the shape of the XRD
patterns. At TA ¼ 823 K the characteristic (100) peak of the B2-RuAl
phase appears at about 29.7 2Q. The shape of the broad diffraction
peak at about 42 suggests, however, that at 823 K as well as at
Fig. 10. DSC trace of the ML with L ¼ 22.44 nm. The constant heating rate was 20 K/
min.
N. Zotov et al. / Intermetallics 18 (2010) 1507e1516
6000
3500
RuAl(110)
RuAl(110)
5000
RuAl(100)
873 K
2000
823 K
1500
698 K
1000
598 K
ML
500
Intensity (arb. units)
Intensity (arb. units)
3000
2500
1513
4000
RuAl (100)
Ru(0002)
Al(111)
3000
873 K
Ru(10-11)
2000
823
1000
698
598
473 K
30
35
40
45
0
25
50
2Θ (degrees)
873 K the MLs have not yet completely transformed. Indeed, XRD
patterns taken at RT on samples cooled from 873 K show the
presence of both ordered B2-RuAl phase and Ru(Al) solid solution.
The in-situ HT XRD patterns of the MLs with L ¼ 2.24 nm show the
same structural behavior with TA.
As has been shown in Section 3.1, MLs with L < 4.5 nm can be
modeled as epitaxially grown MLs. The HT XRD patterns at 473 and
598 K can also be modeled as MLs. However, comparison of the
structural parameters at RT with that at higher temperatures
shows, that the thickness of the Al layers decreases, while the
thickness of the Ru layers increases. Simultaneously, an increase of
the random variation in the Al d-spacings is observed (e.g. from
0.02 nm in the as-deposited ML with L ¼ 2.24 nm to 0.05 nm at
473 K and to 0.08 nm at 598 K). These results suggest that Ru/Al
interdiffusion takes place already at TA 598 K. Indeed, the crosssectional TEM images (Fig. 7) indicate that Ru/Al interdiffusion is
enhanced by Ru/Al intermixing already during deposition.
3.4.2. Multilayers with L > 4.5 nm
Fig. 12 shows the HT XRD patterns for the ML with
L ¼ 22.44 nm. A decrease of the intensity of the peak at about 38.5
2Q is observed up to TA ¼ 598 K, while the Ru (0002) and (10e11)
peaks remain practically unchanged. This gives further support for
the assignment of the peak at about 38.5 2Q as Al (111). At 698 K,
the Al (111) peak disappears completely and some broadening of
the Ru (0002) and Ru (10e11) peaks is observed. At TA ¼ 823 K the
Ru peaks disappear and the characteristic (100) and (110) peaks of
the B2-RuAl phase appear at about 29.7 and 42.5 2Q, respectively.
The in-situ HT XRD patterns of the sample with L ¼ 11.22 nm show
the same structural behavior with TA. The XRD pattern of a sample
cooled from 873 K to RT (Fig. 13) contains only peaks of the B2-RuAl
phase. The calculated lattice parameter of the B2-RuAl phase
a0 ¼ 0.298 0.003 nm is similar to the lattice parameter of bulk B2RuAl (a0 ¼ 0.2996 nm [1, 32]). This shows that the MLs with
L > 4.5 nm transform completely into the B2 phase in contrast to
the MLs with L < 4.5 nm (see Section 3.4.1).
It may be concluded that the first exothermal DSC peak at
about 573 K is due to the mixing of Ru and Al and the formation of
Ru(Al) solid solution, while the second DSC peak corresponds to
the formation of the ordered B2-RuAl phase. We cannot exclude
completely that the first exothermic DSC peak is partially due to
an initial nucleation and lateral growth of B2-RuAl nanograins like
in Ni/Al MLs [33]. However, the volume fraction of such nanograins should be below the detection limit of the XRD technique
used (1e2%).
30
35
40
45
50
55
2Θ (degrees)
Fig. 12. High-temperature X-ray diffraction patterns of the ML with L ¼ 22.44 nm.
The effects of L on the reaction behavior of the investigated Ru/
Al MLs are in agreement with the results for Ni/Al [15,34] and Ti/Al
[16] MLs. In Ni/Al and Ti/Al MLs with small L periods and compositions close to the equilibrium intermetallic phases, the interdiffusion region is comparable to L. Annealing of such MLs leads
directly to the formation of the equilibrium phases (NiAl or TiAl,
respectively) [15,16,34]. For larger L periods, the formation of
intermediate intermetallic phases (h-Al9Ni2, Al3Ni, Ni2Al3 or Al3Ti)
is first observed [15,16,34]. This behavior is similar to the behavior
of the investigated Ru/Al MLs. The phase formation during
annealing of Ru/Al MLs with L ¼ 88 nm starts with the formation of
the intermediate phase RuAl6 [14]. Only at higher temperatures the
transformation to the equilibrium B2-RuAl phase takes place, in
contrast to the present results for Ru/Al MLs with L 22.44 nm.
3.5. Kinetics
Isothermal XRD experiments at 573, 748, 763 and 773 K were
performed on samples with L ¼ 22.44 nm, in order to determine
the kinetics of the two main processes e the formation of the Ru(Al)
solid solution at lower and the B2-RuAl phase at higher temperatures. XRD patterns were measured with 5 min accumulation time.
Fig. 14a and b shows the time dependences of the integral intensity
of the Al (111) peak at 573 K and of the B2-RuAl (100) peak at 763 K,
respectively. The decrease of the intensity of the Al (111) peak
suggests that Ru/Al interdiffusion takes place already at relatively
5000
(110)
Intensity (arb. units)
Fig. 11. High-temperature X-ray diffraction patterns of the ML with L ¼ 4.48 nm.
473
RT
B2 RuAl
4000
3000
2000
(111)
1000
0
35
40
45
50
55
(220)
60
65
2Θ (degrees)
Fig. 13. X-ray diffraction pattern of ML with L ¼ 22.44 nm cooled down from 873 K.
N. Zotov et al. / Intermetallics 18 (2010) 1507e1516
175
150
1200
(RT)
Integral Intensity (arb. units)
a
Integral Intensity (arb. units)
1514
125
100
75
50
Λ = 22.44 nm
TA = 573 K
25
0
5
10
15
20
25
30
35
Annealing Time tA (minutes)
Integral Intensity (arb. units)
600
400
200
690
720
750
780
810
840
870
900
Annealing Temperature TA (K)
Fig. 15. Temperature dependence of the integral intensities of the B2-RuAl (100) peak
for different MLs.
300
demonstrate that the [Ru 1 nm/Al 1.24 nm]80 and [Ru 2 nm/Al
2.48 nm]40 MLs grow almost like ideal MLs perpendicular to the
film surface. The presence of in-plane GBs (see Fig. 5a) could in
principle contribute to the diffusion in these MLs. However, it has
been suggested that for L Lk, in-plane GB diffusion may not be
expected to make a significant contribution [40], where Lk is the
average in-plane grain diameter (Lk w hAki½). Fig. 6a shows that for
MLs with L < 4.48 nm this condition is fulfilled. Therefore, in a first
approximation, the inderdiffusion in these MLs should be governed
mainly by the corresponding bulk interdiffsusion coefficient DML.
But, according to the theory of Cahn and Hiliard [41,42], the interdiffusion coefficient of MLs DML is L-dependent:
250
200
150
100
Λ = 22.44 nm
TA = 763 K
50
0
800
0
0
b
Λ = 2.24 nm
Λ = 4.48 nm
Λ = 11.22 nm
Λ = 22.44 nm
1000
0
20
40
60
80
100
Annealing Time tA (minutes)
Fig. 14. (a) Annealing time dependence of the integral intensity of the Al (111) peak for
the ML with L ¼ 22.44 nm. The dashed line is only a guide for the eye. (b) Annealing
time dependence of the integral intensity of the B2-RuAl (100) peak for the ML with
L ¼ 22.44 nm from isothermal measurements at 763 K. The dashed line is a polynomial
fit.
low temperatures. This confirms the conclusion drawn from the
simulations of the superlattice satellites at 478 and 598 K. On the
other hand, the integral intensity of the RuAl (100) peak increases
with time. Due to the short measuring time, the standard deviations are relatively large. Nevertheless, the intensities can be well
fitted using the parabolic power law I ¼ (c $ tA)0.5, where tA denotes
the annealing time. Similar results were obtained at 748 and 773 K.
Since integral intensities are proportional to volume fraction,
a diffusion-controlled rather than reaction-controlled growth [35]
of the B2-RuAl phase can be assumed. Diffusion-controlled
growth was established also in other aluminide systems [36,37].
To study the influence of L on the kinetics of B2-RuAl formation,
the integral intensity of the B2-RuAl (100) peak, I(100), as a function
of TA is plotted in Fig. 15. Generally, I(100) increases with TA for all
samples. Additionally, an increase of I(100) with increasing L is
observed for each TA. Hence, the kinetics of the nucleation and
growth of the B2-RuAl phase accelerates with L of the Ru/Al MLs.
In ideal (epitaxial) MLs interdiffusion takes place mainly via bulk
lattice diffusion across the interfaces and such metallic MLs have
been widely used for the determination of bulk interdiffusion
coefficients at relatively low temperatures [38]. Recently, Zotov
et al. [39] have shown that vacancy diffusion is probably the main
diffusion mechanism. The XRD results and TEM micrographs
DML(L) ¼ Do $ {[1 þ 2h2 $ Y/f 00] þ 2 $ (k/f 00) $ B2(L)}
(1)
where B2(L) is an orientation-dependent lattice function [43]:
B2(L) ¼ (2/d2o) $ [1 cos(2pdo/L)] z 4p2/L2
(2)
which decreases with increasing L. Do is the true interdiffusion
coefficient, do is the average lattice spacing; h is the logarithmic
derivative of the lattice parameter of the homogeneous phase with
respect of the concentration c; Y is the biaxial modulus of the
material along the growth direction, which for isotropic materials is
equal to E/(1 n), where E is the Young’s modulus and n is the
Poisson’s ratio; f 00 is the second derivative of the Helmholtz free
energy f with respect to c; and k is the coefficient of the first term in
the expansion of f(c) with respect to c [41,42]. The term 2(k/f 00)B2(L)
in Eq. (1) describes the thermodynamic effect of the presence of
concentration gradients at the interfaces. It increases in importance
with decreasing L, because B2(L) w 4p2/L2. This could be qualitatively understood also by the fact that with decreasing L, the
volume fraction of the interfaces increases.
Eqs. (1) and (2) show that the interdiffusion coefficient DML can
be smaller than Do if k/f 00 < 0 or larger than Do if k/f 00 > 0, respectively. For binary systems with negative heat of mixing like CuePd,
CueAu, CoePt etc, k/f 00 is less than zero, because f 00 > 0 and k < 0
[44e46]. Several ordered intermetallics are formed in the Ru/Al
system [1]. Therefore, it might be expected that k/f 00 < 0 for the Ru/
Al system as well. In this case according to Eqs. (1) and (2), the
interdiffusion coefficient DML(L) in the Ru/Al MLs should increase
with increasing L.
As has been shown by XRD, the MLs with L > 4.5 nm are polycrystalline. TEM micrographs show the presence in these MLs of
vertical high-angle GBs and discontinuous atomic layers, seperated
N. Zotov et al. / Intermetallics 18 (2010) 1507e1516
by vertical bridges and island-like cusps (Figs. 7 and 8). These
bridges have lateral lengths larger than the thickness of the interfaces between the Al and Ru atomic layers. Similar to GBs diffusion
in polycrystalline materials [47,48], these microstructural features
in the investigated MLs can serve as fast diffusion pathways (FDP).
At relatively low annealing temperatures TA < ½Tmelting, the diffusivity DFDP along such FDPs is likely to be much larger than the bulk
interdiffusion coefficient (DFDP [ DML) due to lattice disorder,
elevated concentration of vacancies and stronger stress gradients.
Indeed, theoretical investigations of the thermal stability of MLs
[49,50], showed that the ML decomposition is assisted mainly by
rapid mass transport along GBs caused by lateral inhomogeneities
in the individual layers. The inhomogeneities discussed in Refs.
[49,50] e thinning of one of the atomic layers and formation of
bridges between the neighboring second type of layers, are similar
to the microstructural features observed in the investigated MLs.
Experimentally, abnormal fast diffusion along vertical GBs has been
observed in Ti/Al MLs [51]. Direct experimental determination of
the activation energy of GB diffusion in polycrystalline Cu/Ni and
Ta/NiFe MLs confirms that it is much lower than the bulk activation
energy for interdiffusion in MLs [52,53]. The presence of intermixing between the Ru and Al layers during deposition could also
facilitate the formation of the Ru(Al) solid solution and then the
RuAl phase. The TEM micrographs show that the fraction of intermixed regions increases with increasing L.
Therefore, the effective diffusion coefficient Deff of the MLs can
be written, in analogy with the Hart’s equation [47] for GB diffusion
(assuming ‘Type A’ GB diffusion kinetics):
Deff(L) ¼ [1 g(L)]DML(L) þ g(L)DFDP
(3)
where DML(L) is the bulk ML diffusivity given by Eq. (1), DFDP is
the effective diffusivity of the FDPs and g is the fraction of FDPs and
intermixed regions, which most generally will depend on L.
Quantitative determination of g would require performance of
high-resolution three-dimensional X-ray tomography, which is not
a trivial task for thin film samples and is beyond the scope of the
present study. However, the SEM and TEM micrographs, the fraction of the intergranular space, determined by SEM, and the
roughness of the atomic layers, determined by TEM, represent in
our view complementary quantitative measures of g. Figs. 6a and 9
show that these microstructural quantities increase with increasing
L, justifying the conclusion that g increases with L for the investigated MLs. Therefore, the contribution of the first term on the
right side of Eq. (3) will decrease, while the contribution of the
second term will increase. Since DFDP > DML, the effective interdiffusion coefficient of the MLs should increase, due to the increase of
both g and DML with L. The XRD results (Fig. 14b) indicate a diffusion-controlled growth for the B2-RuAl phase. The increase of the
effective interdiffusion coefficient with L, discussed above, means
that the kinetics of the B2-RuAl grain growth should accelerate
with L, as observed (Fig. 15).
The grain sizes of the B2-RuAl phase, determined from the
broadening of the (100) RuAl peak at 823 K, confirm also this
scenario. LRuAl
100 is equal to 10.1 2 nm for the MLs with L < 4.5 nm
and is equal to 27.2 0.4 nm for the MLs with L > 4.5 nm, suggesting faster kinetics in the later case.
1515
A critical period Lc, which indicates a change in the microstructure of the as-deposited MLs and the subsequent B2-RuAl
formation, of about 4.5 nm could be established. Below Lc the Ru
atomic layers grow predominantly along the (0001) direction and
a coherent epitaxy between the Al (111) and the Ru (0001) layers is
observed. The roughness of the atomic layers is small and only
a few structural defects are observed by TEM. Above Lc the epitaxy
is lost. The out-of-plane microstructure of the as-deposited MLs is
polycrystalline with strong Ru (0001) texture. In each Ru atomic
layer there are grains growing along other crystallographic directions, besides the Ru (0001) direction. The individual atomic layers
are significantly distorted and the ML roughness increases with L.
The in-plane microstructure of the MLs is also polycrystalline, as
observed by SEM. The average in-plane grain size of the asdeposited MLs and the fraction of intergranular space increase
with L.
The DSC and HT XRD results show that the MLs transform upon
annealing into the ordered B2-RuAl phase in two stages. Below
w698 K, Al diffusion into the Ru layers takes place, which leads to
the formation of disordered Ru(Al) solid solution. In the second
stage (above w698 K) the ordered B2-RuAl phase is formed. We do
not observe the formation of intermediate intermetallic phases as
in Ru/Al, Ni/Al or Ti/Al MLs with large L [14e16,34]. The nucleation
and growth of the B2-RuAl phase is diffusion-controlled rather than
reaction-controlled.
The kinetics of the B2-RuAl grain growth is L-dependent due to
the expected L-dependent interdiffusion coefficient of the Ru/Al
MLs and the increase of the fraction of FDPs as well as intermixed
regions with L. The kinetics is slower in the as-deposited MLs with
L < 4.5 nm. As a result the transformation of these MLs into the B2RuAl phase is not complete for short annealing times (10 min) at
873 K and the grain sizes of the B2-RuAl phase are rather small
(w10 nm). The MLs with L > 4.5 nm, on the other hand, are polycrystalline and fast interdiffusion along vertical GBs accelerates the
grain growth. Correspondingly, the transformation of these MLs
into the B2-RuAl phase at 873 K even for short annealing times is
complete. Thus, these MLs are most promising for future
applications.
The present results give also further evidence for the presence of
strong L effect on the formation and growth of B2 aluminides from
M/Al (M ¼ Ti, Ni, Ru) multilayers.
Acknowledgement
NZ acknowledges the experimental support of Caesar (Bonn)
and DFG Project EG 101/15-1 for part of these studies.
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