ARTICLE IN PRESS
Vacuum 79 (2005) 171–177
www.elsevier.com/locate/vacuum
Annealing of niobium coatings deposited on graphite
S. Barzilaia,b, A. Raveha,b,, N. Fragea
a
Department of Material Engineering, Ben-Gurion University, P.O. Box 653, Beer-Sheva 84105, Israel
b
Division of Chemistry, NRC-Negev, P.O. Box 9001, Beer-Sheva 84190, Israel
Received 14 December 2004; received in revised form 16 January 2005; accepted 28 March 2005
Abstract
Niobium coatings 8–12 mm thick were deposited by magnetron sputtering on ATJ graphite substrate. The kinetic
growth of carbide layers into niobium coatings and the properties of the Nb2C and NbC phases obtained after
annealing in the temperature range of 1073–1773 K were studied. It was found that the carbide layer growth displayed
parabolic behavior patterns establishing the growth rate constants (K) of Nb2C and NbC layers, as follows:
190 kJ m2
164 kJ m2
K Nb2 C ¼ 3:7 108 exp
; K NbC ¼ 4:5 109 exp
.
RT
RT
s
s
r 2005 Elsevier Ltd. All rights reserved.
Keywords: Sputtering; Coating; Annealing; Niobium carbide
1. Introduction
The surface of commercial graphite can be
improved by the formation of carbide coatings
due to their chemical inertness, hardness and wear
resistance. Niobium carbide has been well documented as an important technological material
because of its unusual combination of physical and
mechanical properties [1]. Therefore, NbC is of
Corresponding author. Division of Chemistry, NRC-Negev,
P.O.
Box 9001, Beer-Sheva 84190, Israel.
Tel.:
+972 544 757 915; fax: +97 286 568 686.
E-mail address: aviraveh@hotmail.com (A. Raveh).
great importance in a wide range of applications,
in corrosive, erosive and wear environments. In
this study, we fabricated niobium-carbide coatings
on graphite substrates by the deposition of
niobium metal followed by thermal annealing in
order to improve the coating to graphite adhesion
and surface properties. In a previous investigation
[2], we showed that the annealing of niobium
coatings produces niobium carbide phases, which
improve the hardness and reduce the porosity of
the coating on graphite. The density, morphology,
and structure of the niobium layer formed by
magnetron radio-frequency (rf) sputtering and of
the carbide derived by thermal annealing are
0042-207X/$ - see front matter r 2005 Elsevier Ltd. All rights reserved.
doi:10.1016/j.vacuum.2005.03.005
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S. Barzilai et al. / Vacuum 79 (2005) 171–177
172
strongly affected by a negative bias voltage (Vb). It
was observed that three distinct categories of Nb
and NbC coatings were formed depending on the
negative bias voltage, Vb: (a) at V b o50 VDC, a
singularly nucleated columnar structure of the Nb
coating was formed, which was transformed into a
highly porous NbC coating by heat treatment; (b)
at 50oV b o80 VDC, a singularly nucleated
columnar structure composed of a continuously
nucleated sub-columnar structure was formed,
which was transformed into a dense NbC coating
with the highest micro-hardness, of 13 GPa; and
(c) at V b X80 VDC, an imperfect structure of the
Nb coating was formed, which was transformed
into a NbC coating with the highest density and an
intermediate range of micro-hardness.
Several authors [3–7] studied the interaction
between Nb and graphite in the temperature range
of 1673–2573 K and they determined the carbon
inter-diffusion coefficients as Nb2C and NbC
layers. Woodford and Chang [3] observed that
the intrinsic diffusivity of carbon into NbC1x is
reduced by the carbide stoichiometry, and that the
diffusion coefficient of niobium is negligible
compared to that of carbon. They also found that
the intrinsic diffusivity of carbon into NbC1x was
sensitive to the phase composition rather than the
activation energy. Brizes et al. [4] demonstrated
that niobium atom diffusion into the carbides is
negligible compared to carbon atom diffusion into
the NbC phase.
The data summarized in Table 1 display the
niobium-carbide layer growth by the diffusion of
carbon into a niobium metal bulk in the temperature range of 1673–2573 K. It summarizes the
kinetic parameters of the interaction between bulk
niobium and carbon, i.e., the constant growth rate
represented by an Arrhenius-type equation, as
shown by other authors [3–7].
Miyake et al. [8] reported the effect of the
graphite substrate temperature on the deposition
of Nb coatings produced by chemical vapor
deposition (CVD). They observed that the coating
deposited at 1473 K was composed of two layers,
identified as Nb and Nb2C phases, while that
deposited at 1523 K was composed of Nb2C and
NbC layers. However, only deposition above
1563 K produced a single NbC layer. They
concluded that the rate of carbon diffusion and
carbide layer formation, originating from a
niobium coating on graphite, is significantly higher
than that observed by other authors [3–6] for
niobium metal bulk.
Results similar to Miyake et al. [8] were
obtained by Isobe et al. [9] while investigating
the interaction between molybdenum coatings on
a graphite substrate. Isobe et al. [9] studied carbon
diffusivity in molybdenum carbide by measuring
the carbide layer growth in both CVD molybdenum coating on a graphite substrate and in bulk
molybdenum that was coupled directly to graphite. They concluded that the diffusion coefficient
for coatings were higher than those obtained for
the bulk specimen. Moreover, the activation
energy was lower than those obtained for the bulk
specimen.
This study was motivated by the lack of data
regarding the growth of carbides fabricated from
niobium coatings. We therefore investigated the
interaction between niobium and graphite by
Table 1
The growth rate constants and carbon diffusivity into niobium carbide bulk as determined by the layer growth
Carbide phase
K0 (m2/s)
QK (kJ/mol)
T (K)
Reference
Nb2C
NbC1x
Nb2C
NbC
Nb2C
NbC
Nb2C
NbC
1.5770.35 105
2.6570.25 105
3 105
1.76 103
2.2 106
4.5 106
5.970.5 107
1.170.1 104
302.5725.8
312.5711.6
337.4
402.2
287.7
305.3
260.1714.1
344.979.1
1673–973
1673–1973
1973–2573
1973–2573
1700–2090
1700–2090
2173–2573
2173–2573
[3]
[3]
[4]
[4]
[5]
[5]
[6]
[6]
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S. Barzilai et al. / Vacuum 79 (2005) 171–177
measuring the carbide layers by means of metallographic and depth profile analyses. The metallographic cross-sections enabled us to distinguish
between sub-layers and the depth profiles verified
the sub-layer composition. These were used to
evaluate the growth rate of the carbide layers
into the niobium coatings formed by thermal
annealing.
173
diluted solutions containing nitric acid and hydrofluoric acid. The thickness of the NbC and Nb2C
layers, obtained after heat treatment of the
niobium coating, were measured by SEM. Measurement of the thicknesses of the sub-layers was
enabled by defining the location of interfaces after
the chemical etching. These measurements were
also verified by depth profile analyses using
wavelength dispersive spectroscopy (WDS), and
micro-probe analyzer.
2. Experimental
2.1. Layer deposition and thermal annealing
3. Results and discussion
The niobium coatings 8–12 mm thick were deposited in a custom-designed rf magnetron system.
The ATJ graphite substrates (10 40 1.5 mm)
were polished with a 600 mesh SiC paper, then
ultrasonically degreased, cleaned and mounted on
a substrate holder. The sputter Nb target (80 mm
in diameter and 6 mm thick, 99.9%) was mounted
on an rf magnetron source at a distance of 5 cm
from the substrate holder. The base pressure was
below 6.67 104 Pa and the sputtering process
was performed in argon (99.999%) at a constant
pressure of 0.67 Pa and an rf input power of
400 W. The substrate holder was subjected to V b ¼
80 V with reference to ground. The maximum
substrate temperature resulting from the deposition process was about 473 K. After deposition,
the Nb layers were annealed in a vacuum of
6.67 104 Pa and in the temperature range of
1073–1773 K for 12–480 min. The heating rate was
3001/min and the cooling rate was 2001/min.
3.1. Nb coatings on graphite
2.2. Layer characterization
Phase analysis was carried out using an X-ray
diffractometer (XRD) with Cu–Ka radiation ðl ¼
0:154 nmÞ and a graphite monochromator. Scanning electron microscopy (SEM) micro-graphs of
the cross-sections after fracturing the sample by
bending tests and also that of metallographic
sections, which were prepared by standard polishing of the samples at several stages of 600, 1000
and 2400 mesh followed by polishing with a
synthetic cloth and 1 mm diamond paste were
taken. The metallographic samples were etched in
The structure and properties of the Nb coatings
were studied as a function of gas pressure, rf
power, negative bias voltage and deposition time.
It was found that the bias voltage and gas pressure
were the two parameters that mainly affect the
micro-structure and the deposition rate [2]. The
change in the micro-structure was seen from
singularly nucleated columnar structure to continuous nucleated structure. In addition, the bias
voltage also causes the nucleation of sub-columnar
structure, while coatings deposited at V b X80 V
show broken columnar structure with a smoother
surface. The effect of the pressure and bias voltage
were seen to be in qualitative agreement with the
structure zone T of Thoronton diagram [10,11].
The coatings deposited at various Vb and at
5 mTorr contained micro-pores [2]. Three ranges
of micro-porosity were observed: (a) at
Vb ¼ 0–50 V, the coatings show 10 vol%; (b) at
Vb ¼ 50–80 V, the coating show 5–8 vol%; while
coatings deposited at range (c) V b 480 V showed
2–3 vol% of micro-porosity. It is probable that the
micro-structure and the micro-porosity concentration affect the carbide growth rate and the final
structure after annealing.
Fig. 1 depicts the fracture of the Nb layers
deposited at various bias voltages. The micrographs represent the three categories of Nb coatings
which were formed by the various negative bias
voltages, Vb: (a) at V b ¼ 0 V, singularly nucleated
columnar (SNC) structures with relatively smooth
domed tops were formed (Fig. 1a); (b) at Vb of
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174
S. Barzilai et al. / Vacuum 79 (2005) 171–177
became smoother (Fig. 1c). The differences in the
columnar structures are attributed to the increased
intensity of ion bombardment with the increase in
negative bias voltage.
3.2. Nb– C layer growth rate
XRD peak patterns of the deposited 8–12 mm
thick coatings were examined after annealing for
3 h at various temperatures (Fig. 2). In thin 2 mm
coatings treated at low T ¼ 1173 K, X-ray depth
sensing could only identify a Nb2C phase, but
could not identify it in thick coatings (curve b). At
T ¼ 1373 K, a Nb2C phase was formed (curve c),
but only after treatment at T ¼ 1773 K did the
treated Nb layer transform to a single NbC phase
(curve d). It is probable that Nb2C and NbC
phases were formed in the interface between the
coating and graphite at 1173 and at 1373 K,
respectively. However, these phases could not be
detected in thick coatings, X5 mm, due to the
XRD method of depth sensing.
Metallographic cross-section micro-graphs enabled us to distinguish between the phases by
displaying clear boundaries, as can be seen in
Nb
700k
Nb2C
Intensity [arbitrary units]
600k
Fig. 1. Effect of the negative bias voltage on the fracture
morphology of Nb coatings deposited on graphite: (a) V b ¼ 0;
(b) V b ¼ 50 V; and (c) V b ¼ 80 V.
500k
NbC
(d)
1773 K
(c)
1373 K
(b)
1173 K
(a)
as deposited
400k
300k
200k
100k
0k
50 V, secondary nucleation occurred, causing the
columnar structure to transform into a continuously nucleated sub-columnar (CNSC) structure
with a rough surface (Fig. 1b); (c) at V b ¼ 80 VDC,
the columnar structure broke down and the surface
30
35
40
45
2θ [degrees]
Fig. 2. Typical XRD patterns of Nb coatings: (a) after
deposition; (b) after 3 h heat treatment at temperatures of
1173 K; (c) 1373 K; and (d) 1773 K.
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S. Barzilai et al. / Vacuum 79 (2005) 171–177
0
67
133
t [min]
200 267 333 400
467 533
1073 K
1173 K
1273 K
1373 K
1473 K
2.0
W2Nb2C*1011 [m2]
Fig. 3. The metallographic section after chemical
etching revealed the boundaries caused by a
preferred chemical attack between the phases. In
addition, fractured SEM micro-graphs (not
shown) also demonstrated distinct Nb, Nb2C and
NbC phase structures.
Fig. 4 depicts the square thickness (W2) of the
NbC and Nb2C phases (Fig. 4) versus duration
time of treatment t at various temperatures T. A
linear relationship was observed using the least
square method which enabled us to define the
growth rate constant (K) at various T. By plotting
the ln K vs 1/T for the NbC and Nb2C phases
(Fig. 5), Arrhenius-type relationships were obtained. It was observed that the K values are
expressed by Eqs. (1) and (2), as follows:
190 kJ m2
8
K Nb2 C ¼ 3:7 10 exp
,
(1)
RT
s
175
1.5
1.0
0.5
0.0
0
4
8
12
(a)
0
83
167
16
20
t*10-3 [s]
24
t [min]
250
333
28
417
32
500
1.4
1073 K
1173 K
1273 K
1373 K
1473 K
1573 K
1673 K
1773 K
W2NbC*1010 [m2]
1.2
1.0
0.8
0.6
0.4
0.2
0.0
0
(b)
5
10
15
20
25
30
t*10-3 [s]
Fig. 4. The square thicknesses (W2) vs heat treatment duration:
(a) for Nb2C layer at 1073–1273 K; (b) for NbC layer at
1073–1773 K.
164 kJ m2
K NbC ¼ 4:5 109 exp
,
RT
s
Fig. 3. Metallographic and fractured cross-section of niobium
coating: (a) after treatment at 1273 K for 12 min; (b) after
treatment at 1273 K for 100 min.
(2)
where K Nb2 C and KNbC are the growth rate
constants of Nb2C at T ¼ 1073–1473 K and NbC
at T ¼ 1073–1773 K. The calculated values of
K Nb2 C and KNbC are given in Table 2.
For comparison, K values for carbon in bulk
niobium obtained by other authors [3–6] are
plotted in Fig. 5. It was observed that the Q and
K values obtained in our study for niobium thin
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176
1E-11
Resnick [6]
287.7 kJ/mole
KNb2C [m2/s]
1E-12
Woodford [3]
302 kJ/mole
1E-13
1E-14
Brizes [4]
337.5 kJ/mole
1E-15
Bornstein [5]
259.7 kJ/mole
this work
190 kJ/mole
1E-16
1E-17
4
5
6
8
9
10
Resnick [6]
344.6 kJ/mole
1E-11
Woodford [3]
312.4 kJ/mole
1E-12
KNbC [m2/s]
7
104/T [K-1]
(a)
1E-13
1E-14
Brizes [4]
402.3 kJ/mole
1E-15
this work
164.2 kJ/mole
Bornstein [5]
305.3 kJ/mole
1E-16
4
(b)
5
6
7
8
9
10
104/T [K-1]
Fig. 5. The growth rate constant (K) for: (a) Nb2C phase; and
(b) NbC phase.
Table 2
The growth rate constants of NbC and Nb2C coatings
Temperature (K)
KNb2C (m2/s)
KNbC (m2/s)
1073
1173
1273
1373
1473
1573
1673
1773
3.4 1017
2.5 1016
7.8 1016
2.8 1015
7.3 1015
1.7 1014
7.8 1014
3.1 1014
2.6 1017
7.7 1017
4.6 1016
2.8 1015
6.1 1015
—
—
—
—, not examined.
coatings differ significantly from those obtained
for bulk niobium, i.e., Q is 60% and K is 3–7
times greater than those values obtained for bulk
niobium. Similar differences in the Q values
obtained for thin coatings versus bulk molybdenum were found by Isobe et al. [9]. They found
that the Q value for the formation of a carbide
phase from a thin molybdenum layer is in the
range of 138–180 kJ/mol depending on the deposition process, in comparison to values reported for
bulk material which varied between 319 and
380 kJ/mol [12–14]. It is worth mentioning that
the data presented in Fig. 5 were taken from
different temperature ranges. According to
Matzke [15,16], the activation energies can vary
with varying temperatures due to the presence of
different diffusion mechanisms, as will be discussed below. In fact, the difference between the
kinetic carbide growth rate in a thin metal layer
compared with bulk material is well documented
[9,13,14,17]. However, these differences are not
fully explained and apparently not well understood. We believe that it may be caused by the
micro-porosity found in the thin layer [2]. It is
possible that the diffusion rate onto the micropore surface in the coating is higher than the one
which occurs in the already-formed carbide phase.
In addition, it is also possible that the activation
energy through the layer surface is lower compared to that through bulk.
The growth rate constants of NbC and Nb2C
are dependent on carbon diffusion into niobium
carbide. Possible reasons why the thin layer
diffusivities differed from those of bulk can be
related to high defect concentration such as that
caused by micro-pores, dislocations induced by the
energetic bombardment, and high quenching rates
inherent in the deposition technique. Microporosity causes surface diffusion on the interface
between the pores and can generate an enhanced
diffusion path for carbon. This may cause the
increase of the carbide layer growth rate predominantly at low temperature interdiffusion ranges,
i.e., 0.35 of the melting temperature. In this
temperature region, short-circuiting processes,
such as surface and dislocation diffusion, are
dominant mechanisms contributing to carbide
growth rate.
Carbon atom activation energy (representing
the migration enthalpy) is necessary for carbon
atom movement from one sub-lattice site to
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another [18,19]. Migration enthalpy is highly
dependent on atom proximity. While carbon
atoms in the lattice are entirely surrounded by
other atoms, those on the surface or in the
dislocation are not. This facilitates atom migration
in the presence of any lattice defects. For example,
in FCC structures it is well known that the surface
activation energy is about 2.6 times lower than
that in bulk diffusion.
In conclusion, thin layers diffusivities differed
from those of bulk because of micro-porosity
and/or high dislocation density inherent in the
deposition technique, enabling atom carbon migration in surface, dislocation and bulk diffusion. This raises the diffusion coefficients (mainly
at low temperature) and lowers the activation
energy of thin layers in contrast to bulk diffusion
processes.
4. Summary and conclusions
Niobium-carbide layers were formed after niobium coatings were deposited on graphite by
magnetron sputtering followed by heat treatment
at various temperatures and time durations. Our
study arrived at the following conclusions:
(a) Structure and micro-porosity are important
factors for determining carbide growth rate
and can explain the difference in the values
found in thin layers versus bulk materials.
(b) The growth of the carbide phases displays
parabolic behavior enabling us to determine
the growth rate constants of the NbC and
Nb2C sub-layers.
(c) The activation energies of the growth of the
NbC and Nb2C sub-layers were 190 and
164.2 kJ/mol, respectively.
177
Acknowledgments
The authors wish to thank Mr. Avi Ben-Shabat
for his expert technical assistance, and Mr. E.
Boublil for SEM micro-graphs. This work was
supported by a grant from the Israeli Council of
High Education and the Israeli Atomic Energy
Commission.
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