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Journal of Nuclear Materials 441 (2013) 372–379
Contents lists available at SciVerse ScienceDirect
Journal of Nuclear Materials
journal homepage: www.elsevier.com/locate/jnucmat
Effect of temperature on diffusion behavior of Te into nickel
Yanyan Jia, Hongwei Cheng, Jie Qiu, Fenfen Han, Yang Zou ⇑, Zhijun Li ⇑, Xingtai Zhou, Hongjie Xu ⇑
Thorium Molten Salts Reactor Center, Shanghai Institute of Applied Physics, Chinese Academy of Sciences, Shanghai 201800, China
a r t i c l e
i n f o
Article history:
Received 25 October 2012
Accepted 18 June 2013
Available online 28 June 2013
a b s t r a c t
The diffusion behavior of tellurium into pure nickel is investigated at different annealing temperatures.
The purpose is to understand the diffusion mechanism of Te by studying the influence of Te on the surface
morphology and grain boundary of nickel. The results show that the surface morphology, reaction product and penetration of Te are greatly dependent on annealing temperature. The depth of Te penetration
increases with the elevated temperature from 500 °C to 1000 °C, and the surface reaction products is
NiTe0.67 (Ni3Te2) or mixture of a spot of NiTe0.69. Te diffuses into nickel predominantly along the grain
boundary at low temperature (below 900 °C), while the diffusion mechanism of Te turns to lattice diffusion above 1000 °C.
Crown Copyright Ó 2013 Published by Elsevier B.V. All rights reserved.
1. Introduction
Nuclear power will be the best choice to meet the future energy
demand for the advantages in environment protection, fuel economy, saving cost, etc. Molten-salt reactor (MSR), one of the most
promising next generation reactors, is different from the other five
generation IV reactors. In details, the other fourth generation reactors use solid-fueled rods, while the fuel is dissolved in molten
fluoride salt in MSR [1–3]. MSR has incomparable advantages compared with other solid-fuel-element systems: a high negative temperature coefficient of reactivity, producing less long-lived wastes
by a factor of 10–100, removing fission products online, excellent
heat transfer characteristics, a very low vapor pressure, thermodynamic stability, and so on [2,4]. The concept of the MSR [5,6] was
developed in the 1950s and then a small Molten Salt Reactor
Experiment (MSRE) was successfully designed, constructed and became critical in 1965. Moreover, it had normally operated about 20
thousand hours by Oak Ridge National Laboratory (ORNL) [6,7].
The metal structure components in MSRE were exposed to
several environments, the main of which were high temperature,
severe corrosion of the molten fluoride salts, and neutron irradiation damage. In order to resist those harsh environment conditions,
Hastelloy N was developed specifically by ORNL for using in
molten-salt systems [8]. However, operation of MSRE revealed
some deficiencies in this alloy. One of them was that the grain
boundaries were embrittled to depths of 50–250 lm in all
⇑ Corresponding authors. Address: 2019 Jia Luo Road, Jiading District, Shanghai
Institute of Applied Physics, Chinese Academy of Sciences, Shanghai 201800, China.
Tel.: +86 21 39194767; fax: +86 21 39194769 (Z. Li).
E-mail addresses: [email protected] (Y. Jia), [email protected] (Y. Zou),
[email protected] (Z. Li), [email protected] (H. Xu).
Hastelloy N exposed to the fuel salt, no matter whether it is in
the irradiational environment [9]. Intergranular cracks formed on
the surface of alloy could deteriorate the mechanical properties
of alloy seriously, which could result in catastrophic accident of
nuclear power. Therefore, it is very important to understand the
formation mechanism of intergranular cracks clearly and modify
the alloy. Considerable evidences have indicated that the formation of cracks on surface of Hastelloy N was mainly related to the
inward diffusion of the fission-product tellurium (Te) [6,10]. The
concentration of Te in MSR would increase with the service time
or power of reactor. According to ORNL reports [10,11], it would
produce 1019 atoms/cm2 of Te over a 30-years period in molten salt
reactor, which was amount to 1 mg/cm2 of Te on the metal surface
based on the assumption of uniform dispersion on the metal surface and diffusion to a depth of 5 mils into the surface layer [6].
The crack susceptibility of a wide variety of metals caused by
the penetration of Te was compared in ORNL technical reports
[11]. It was explicitly pointed out that the intergranular crack
depth would increase with the rise of annealing time and temperature, and the content of Te, whereas no intergranular cracks existed in stainless steel and copper. But no detailed explanation
was elaborated in these reports and the role of Te on intergranular
crack formation in nickel based superalloys is unclear so far. Therefore, further study should be performed to confirm the exact damage mechanism of Te.
Fortunately, Te probably behaves in non-fissioning melts much
as it does in a fissioning salt, and intergranular cracks have nothing
to do with the corrosion [8]. So that laboratory experiments can be
used to simulate the environment for crack formation and answer
many questions. The researchers of ORNL had studied the effect of
Te on Hastelloy N and other alloys by electroplating or vapor deposition, and proved that these methods of experiments are feasible
0022-3115/$ - see front matter Crown Copyright Ó 2013 Published by Elsevier B.V. All rights reserved.
http://dx.doi.org/10.1016/j.jnucmat.2013.06.025
Y. Jia et al. / Journal of Nuclear Materials 441 (2013) 372–379
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and reasonable [6–9,11]. Therefore, in the present study, we intent
to research the diffusion mechanism of Te in nickel (Ni) first to reveal the relations between Ni and Te used the electroplating method for preparing samples.
2. Experimental procedure
Fig. 2. Sample encapsulated in a quartz ampoule.
In this work, Te was electroplated on the surface of 99.9% comm
ercial Ni. The specimens with dimensions of 20 mm 10 mm 0.5 mm were cut by electrospark wire-electrode cutting
from the Ni plate. The surfaces of these specimens were prepared by
grinding on metallographic SiC papers to #800, then ultrasonically
cleaned in acetone and rinsed with deionized water before electroplating. The electroplating solution was prepared by dissolving the
telluric oxide in hot one molar potassium hydroxide, about 100 ml/
g Te.
The schematic diagram of electroplating device is shown in
Fig. 1. The platinum foil was used as anode and the electroplated
sample as cathode. The plating potential was 15 V DC and the current was 10 mA/cm2. Te was electroplated and its content was
about 0.5 mg/cm2 on the surface of Ni foil. After electroplating,
all the specimens were washed in deionized water for 20 s and
then put into the drying oven at 75 °C for about 6 min until the surface color of foil was changed into blue-gray. Each specimen sealed
in a vacuum quartz ampoule which pressure was less than 5 Pa to
remove the influence of the oxygen (Fig. 2). Then it was annealed
for 100 h in a muffle furnace at 500 °C, 600 °C, 700 °C, 800 °C,
900 °C and 1000 °C respectively, and then removed from the furnace and cooled in air to examine the microstructure evolution.
Scanning electron microscopy (SEM) equipped with energy dispersive spectroscopy (EDS) and wavelength dispersive spectrometer (WDS), Metallograph, X-ray fluorescence (XRF), Electron probe
microanalysis (EPMA) and X-ray diffraction (XRD) were used to
characterize the sample morphology, Te diffusion depth and to
identify the reaction products. Inductively coupled plasma mass
spectrometry (ICP-MS) was used to measure Te of very low concentration in the sample. The etchant consisting of deionized
water, acetic acid and nitric acid (20:100:38) was selected to etch
the sample for observing the cross-section morphology.
For the XRF analysis, two-dimensional micro-XRF (400 lm 70 lm and 60 lm 500 lm) were acquired at the micro beamline
15U of Shanghai Synchrotron Radiation Facility (SSRF), Shanghai,
China. The spot size is 3.4 lm (horizontal FWHM)3.2 lm (vertical
FWHM), measuring time was 3 min real time and step size was set
to 3.5 lm.
Fig. 1. Schematic diagram of electroplating device: (a) 15 V DC power supply, (b)
sample, and (c) platinum gauze electrode.
Fig. 3. XRD patterns of foils annealed at different temperatures for 100 h.
3. Results
Typical X-ray diffraction as the function of annealing temperature on the diffusion of Te in Ni is shown in Fig. 3. At room temperature, i.e. non-annealing treatment before and after electroplating,
all diffraction peaks were identified to the phases of Ni matrix and
film Te displayed in Fig. 3 RT (no Te) and RT (with Te). In the
annealing temperature regions from 500 to 900 °C, NiTe0.67 (Ni3Te2) perhaps accompanied with a small amount of NiTe0.69 appeared and all the peaks changed with increasing temperature
indistinctively. Nevertheless, only Ni peaks existed after annealing
at 1000 °C.
The surface morphologies of Te-electroplated Ni foils annealed
at different temperatures for 100 h are compared in Fig. 4. The original surface of pure Ni was shown in Fig. 4a. At room temperature
(RT), the needle-like clusters developed on the surface of foil
(Fig. 4b) were elemental Te detected by XRD method (Fig. 3). After
annealing treatment, many particles were closely packed on the
surface of foil as shown in Fig. 4c–e. Although the phase composition was similar at the range temperature of 500–900 °C according
to XRD analysis (Fig. 3), it was worth noting that a few lamellar
phases (as shown by arrows in Fig. 4d and e) formed on the surface
of Te-electroplated nickel foil annealed from 500 °C to 700 °C,
whereas only large blocky particles formed at 800 °C and 900 °C
(Fig. 4f and g). When the annealing temperature was increased further to 1000 °C, the surface morphology of foil changed into terrace-shape which was totally different from that of others.
Fig. 5 shows the cross-section morphologies of Te-electroplated
Ni foil annealed at different temperatures for 100 h. All of them
were etched in the same conditions. The grain boundary morphology of Ni foil varied obviously with annealing temperature after
etching. For the foil non-annealed and annealed at 500 °C, the grain
boundaries near the surface were similar to those of matrix (Fig. 5a
and b). However, they were etched more seriously than inner grain
boundaries with increasing of annealing temperature from 600 °C
to 800 °C. When the sample exposed at 900 °C, the morphology
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Y. Jia et al. / Journal of Nuclear Materials 441 (2013) 372–379
Fig. 4. Surface morphologies of foil annealed at different temperatures for 100 h: (a) non-annealing before Te plating, (b) non-annealing after Te plating, (c) 500 °C, (d) 600 °C,
(e) 700 °C, (f) 800 °C, (g) 900 °C and (h) 1000 °C.
near the grain boundaries and surface were corroded over a wide
range and much darker than the matrix (Fig. 5f), For the foil annealed at 1000 °C, the grain boundary was not so serious corrosion
compared with the foil annealed in the temperature range from
700 °C to 900 °C, but the matrix had darker morphology than that
at 600 °C.
Considering the various grain boundary morphology in Fig. 5,
the content of Te in the grain boundary and matrix close to the surface of Ni foil annealed at 700 °C is analyzed by WDS equipped on
SEM as shown in Fig. 6 and Table 1. It was observed that the depth
of Te penetration into nickel along the grain boundary was little
more than that in the matrix relatively. And the content of Te
was trace amount. (The limit of detection of WDS is 0.05 atomic%
and each data was an average value calculated on three tested
data). That is because the fraction of Te which actually diffuses into
the sample is probably quite small, with the majority reaction to
form nickel telluride on the foil surface [6]. Moreover, its content
is lower than the actual value at the grain boundaries after being
etched (as shown in Table 1).
Fig. 7 illustrates the element distribution of cross-section foil
annealed at different temperatures for 100 h. SEM and EDS results
with the combination of XRD analysis suggested that the loose and
porous reaction layer of Ni and Te formed on the surface of foil annealed at 500–900 °C. According to EDS results obtained from line
scan analysis, a diffusion layer of Te element which thickness changed not obviously (2 lm) exists below the reaction layer in the
range of 500–900 °C. The content of Te decreased dramatically in
the diffusion layer and even almost could not be detected in Ni matrix. Strangely, the cross-sectional morphology and element distribution of foil annealed at 1000 °C was different from that at other
temperatures. No reaction layer was found on the surface of foil,
and Te wasn’t detected by EDS.
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Y. Jia et al. / Journal of Nuclear Materials 441 (2013) 372–379
Fig. 5. Cross-section metallographs of Te-electroplated Ni foil annealed at different temperatures for 100 h: (a) non-annealing after Te plating, (b) 500 °C, (c) 600 °C, (d)
700 °C, (e) 800 °C, (f) 900 °C and (g) 1000 °C.
Fig. 6. SEM erosive cross section images of Te-plating sample annealing at 700 °C
for 100 h. (a) Grain Boundary, and (b) matrix.
Fig. 8 shows the l-XRF analysis results of the sample annealed
at 800 °C and 1000 °C corresponding local area (rectangular area in
cross section). It was clear that the diffusion depth of Te in Ni foil
annealed at 1000 °C was much deeper than that at 800 °C. The fluorescence signal of Te was only about 25 lm depth at 800 °C,
whereas the detected depth was almost 250 lm for the foil annealed at 1000 °C. The mapping (Fig. 8b) also showed that the
intensity of Te in Ni foil annealed at 1000 °C gradually decreased
from surface towards matrix in the external layer. The discrepancy
of diffusion depth of Te in Ni at 800 °C and 1000 °C would be related to the diffusion mechanism of Te.
Table 1
Content of Te corresponding to the point in Fig. 6a and b.
Point
1
2
3
4
5
6
7
Te (atomic%)
Ni (atomic%)
0.92 ± 0.003
99.07 ± 0.007
0.08 ± 0.005
99.91 ± 0.005
0.07 ± 0.002
99.92 ± 0.008
0.04 ± 0.008
99.95 ± 0.002
6.05 ± 0.007
93.94 ± 0.003
0.10 ± 0.009
99.89 ± 0.001
0.02 ± 0.004
99.97 ± 0.006
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Y. Jia et al. / Journal of Nuclear Materials 441 (2013) 372–379
Fig. 7. Cross-sectional SEM images and corresponding EDS line scan of Te-electroplating Ni foil annealed at different temperatures for 100 h: (a) no deposited Te, (b) 500 °C,
(c) 600 °C, (d) 700 °C, (e) 800 °C, (f) 900 °C and (g) 1000 °C.
Y. Jia et al. / Journal of Nuclear Materials 441 (2013) 372–379
377
Fig. 8. Cross-sectional scanning position and corresponding micro-XRF maps of Te in Ni section: (a) 400 lm 70 lm for the sample annealed at 800 °C, and (b)
60 lm 500 lm for the sample annealed at 1000 °C. The graduated color bar from red to blue shows the change of Te counts increment. (For interpretation of the references
to color in this figure legend, the reader is referred to the web version of this article.)
4. Discussion
In the present study, the experimental evidences of reaction and
diffusion of Te in Ni have been obtained. According to the above
experiment results, reaction product and diffusion mechanism of
Te in purity Ni would change with various annealing temperature.
According to Rosenthal and Briggs’s study, the reaction product
between Ni and Te is changed with increasing of temperature [6].
And beta, gamma, and delta phases of nickel telluride were identified in the Ni-Te system by Barstad et al. [12]. Nickel telluride was
detected by XRD analysis from 500 °C to 900 °C may contains one
or two kinds of crystallographic structure of phase in the present
study. But it is difficult to differentiate them specifically by XRD
alone, because many of them have virtually identical interplanar
spacings. It is well known that the beta field extends from NiTe0.67
(Ni3Te2) to NiTe0.7 and has three crystallographic structures, and
Ni3Te2 appears to be the most stable and prevalent nickel telluride
in the range of 600–700 °C [12]. The various morphology of the
reaction products developed on the surface of Ni at different temperatures could account for the transformation process of nickel
tellurides from non-stable state to stable. Therefore, Ni3Te2 is deduced as the main stable nickel telluride annealed at low temperatures in present study.
It has been reported that nickel telluride decomposed into Ni
and Te vapor above 800 °C [6]. But the diffraction peaks of Ni3Te2
are still observed (Fig. 3) and brittle layer appears on the surface
of sample (Fig. 7) when sample is annealed at 900 °C in this study.
It is because the annealing time of nickel telluride is too short to
decompose completely. With annealing temperature increasing
to 1000 °C, no diffraction peaks of nickel telluride are detected by
XRD. Nevertheless, the result of ICP-MS shows that the mass percentage content of Te is about 800 ppm in Ni matrix, which means
that a great majority of Te has diffused into matrix. It is also confirmed by Synchrotron Radiation micro-XRF (Fig. 8b). With the
increasing of temperature, the transformation or decomposition
of nickel telluride will influence on the microstructure of the foil.
From 600 to 900 °C, the morphology of grain boundaries adjacent to the surface of foil seems like coarser than those in center
position. And with the temperature increasing, the phenomenon
is more and more obvious, which might be related to the diffusion
behavior of Te. The grain boundary has higher energy state and
lower diffusion activation energy compared with the matrix [13].
Therefore, Te diffuses easily into nickel along the grain boundary,
and the diffusion behavior is related to temperature [14]. At low
temperature, the diffusion rate is slow and only few of Te diffuses
into Ni along the grain boundary. So the grain boundary is not
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Fig. 9. The element distribution of Te for the sample annealed at (a) 700 °C, (b) 800 °C, (c) 900 °C and (d) 1000 °C. The graduated color bar from red to blue shows the change
of Te counts increment. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
etched easily by metallographic etchant below 600 °C. With the
temperature increasing above 800 °C, the atom diffusion velocity
of Te is accelerated and Te diffuses easily into matrix along the
grain boundary, and the corrosion resistant of grain boundary of
nickel is getting worse. Thus, the grain boundary is vulnerable to
etching by etching solution and its morphology is much coarse as
shown in Fig. 5e and f. While at 1000 °C, the thermal vibration of
the matrix atoms is much more severe because of higher temperature. Therefore, the migration rate of Te is faster and it can diffuse
easily into Ni not only along the grain boundaries but also the matrix at 1000 °C or above. As a result, lattice diffusion becomes obvious and Te distributes equally into matrix [15]. It results in that the
content of Te in grain boundary is lower than that at the other temperatures. So the grain boundary of the sample annealed at 1000 °C
is not etched so serious compared with that at 800 and 900 °C.
The element distributions of different samples are directly measured after polished with electron-microprobe analysis, as shown
in Fig. 9. It can be obviously noted that Te diffused gradually into
nickel along grain boundary below 900 °C predominantly
(Fig. 9a–c), meanwhile accompanied with lattice diffusion in particular 900 °C, whereas the diffusion mold obviously changed into
lattice diffusion at 1000 °C. The result of EPMA is consistent with
the above analysis. According the above analysis, it could be concluded that Te diffuses into nickel along the grain boundary at
low temperature, and the diffusion mode will transfers gradually
from grain boundary to lattice diffusion with increasing of
temperature.
1. With the annealing temperature increasing from 500 to 900 °C,
the reaction products on the surface of Ni is mainly NiTe0.67 or
mixture of a spot of NiTe0.69. The surface morphologies of Teelectroplated Ni foil are changed from cluster-type needles to
small thin platelets. At 1000 °C, its surface morphology turns
into terrace-shape.
2. Te diffuses into Ni predominantly along grain boundary at low
temperature (below 900 °C) and the diffusion is more severe as
the exposure temperature was increased, whereas its mechanism turns into lattice diffusion at 1000 °C.
Acknowledgements
This work was supported by Science and Technology Commission of Shanghai Municipality (Grant No. 11JC1414900), Project
supported by Thorium Molten Salts Reactor Fund (Grant No.
XDA02000000), Project supported by the National Natural Science
Foundation of China (Grant Nos. 11005148 and 50904044), Project
supported by the Special Presidential Foundation of the Chinese
Academy of Science, China (Grant No. 29), and Project supported
by the National Basic Research Program of China (Grant No.
2010CB934501). This research was carried out at Key Lab of Nuclear Analysis and Nuclear Energy Technology, Chinese Academy of
Sciences and the beamline 15U of Shanghai Synchrotron Radiation
Facility (SSRF), China.
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