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Deuterium permeation through erbium oxide coatings on RAFM steels by a dip-coating technique
Journal of Nuclear Materials 442 (2013) 533–537
Contents lists available at SciVerse ScienceDirect
Journal of Nuclear Materials journal homepage: www.elsevier.com/locate/jnucmat
Deuterium permeation through erbium oxide coatings on RAFM steels by a dip-coating technique Takumi Chikada a,⇑, Shunya Naitoh b, Akihiro Suzuki c, Takayuki Terai a, Teruya Tanaka d, Takeo Muroga d a
Institute of Engineering Innovation, School of Engineering, The University of Tokyo, 2-11-16 Yayoi, Bunkyo-ku, Tokyo 113-8656, Japan Department of Systems Innovation, Faculty of Engineering, The University of Tokyo, 7-3-1 Hongo, Bunkyo-ku, Tokyo 113-8656, Japan c Nuclear Professional School, School of Engineering, The University of Tokyo, 2-22 Shirakata-shirane, Tokai, Naka, Ibaraki 319-1188, Japan d Fusion Systems Research Division, Department of Helical Plasma Research, National Institute for Fusion Science, 322-6 Oroshi, Toki, Gifu 509-5292, Japan b
a r t i c l e
i n f o
Article history: Available online 7 June 2013
a b s t r a c t A tritium permeation barrier is a promising solution for the problems of tritium loss and radiological safety in fusion blanket systems. In recent years, erbium oxide coatings have shown remarkable permeation reduction factors. One of the remaining issues for the coatings is the establishment of plant-scale fabrication. In this study, erbium oxide thin films have been fabricated by a dip-coating technique, which has the potential to coat a complex-shaped substrate, and deuterium permeation behavior in the coatings has been examined. Crack-free coatings were formed on a reduced activation ferritic/martensitic steel F82H substrate by use of a withdrawal speed of 1.0–1.4 mm s1 and a heat-treatment process in hydrogen with moisture. In deuterium permeation experiments, a 0.2-lm-thick coating on both sides of the substrate showed a reduction factor of 600–700 in comparison with a F82H substrate below 873 K; however, the coating degraded at above 923 K because of crack formation. A double-coated sample indicated a reduction factor of up to 2000 and did not degrade at up to 923 K. The driving pressure dependence of the deuterium permeation flux indicated that the permeation tended to be limited by surface reactions at low temperatures. Optimization of the number of layers has the possibility to reduce degradation at high temperatures while maintaining high permeation reduction factors. Ó 2013 Elsevier B.V. All rights reserved.
1. Introduction Control of tritium inventory and tritium leakage has gained attention for the development of deuterium–tritium (D–T) fusion reactors. One of the most crucial issues is tritium permeation through structural materials in blanket systems because of fatal fuel loss and radiological hazards. Ceramic coatings have been investigated for several decades to provide a thin hydrogenimpermeable layer on inner pipe walls as a tritium permeation barrier (TPB) [1–4]. In recent years, erbium oxide (Er2O3) coatings have been assessed as being a promising candidate for a TPB [5–7]. A series of studies achieved precise deuterium permeation behaviors through high-purity Er2O3 coatings deposited on reduced activation ferritic/martensitic (RAFM) steel substrates as well as permeation reduction factors (PRFs) of up to 105 at 873 K [8–10]. Recently, the status of the investigation has progressed to the stage of developing a fabrication process toward practical utilization such as plant-scale fabrication without restriction of s-
⇑ Corresponding author. Tel./fax: +81 3 5841 7420. E-mail address: [email protected] (T. Chikada). 0022-3115/$ - see front matter Ó 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jnucmat.2013.05.072
ubstrate geometry. One of the ongoing fabrication methods is metal–organic chemical vapor deposition (MOCVD). Fabrication of the Er2O3 coating on the tubular substrates was demonstrated. However, at present, a 1-lm-thick coating decreased hydrogen permeation to 1/20 of that with a SS316 plate [11]. Another is metal–organic decomposition (MOD), which is a liquid phase method with an Er-contained precursor. The liquid phase method has a great advantage in terms of simplicity of fabrication facilities (e.g., a large vacuum system is unnecessary), but control of coating properties such as crystallinity, the amount of impurities, and intermediate layers is challenging. In our previous study, a 0.3-lm-thick MOD sample coated by a spin-coating technique showed a PFR of 500–700 by controlling oxygen potential during the heat treatment [12]. However, in the case of the MOD samples coated by dip coating, approximately 1-lm-thick coatings on both sides of the RAFM substrate showed PRFs of up to 102 [13,14]. This result is probably due to the formation of cracks or peelings reaching the substrate even if the Er2O3 layer has been sufficiently crystallized. The goal of this study is to optimize the dip-coating parameter for the fabrication of crack-free MOD coatings and to acquire fine hydrogen permeation behaviors in the coatings.
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2. Experimental
3. Results
2.1. Sample preparation
As the first step in coating process optimization, surface profiles of the coated samples withdrawn at 0.05–1.5 mm s1 were compared, as shown in Fig. 1. The samples withdrawn at below 0.75 mm s1 showed spots. The sample withdrawn at 1.5 mm s1 showed obvious unevenness of the coating. No clear peeling was observed in the samples withdrawn at 1.0–1.4 mm s1. In particular smooth surfaces were observed on the samples withdrawn at 1.0–1.2 mm s1. Interfacial textures of the coatings were examined by cross-sectional observation, as shown in Fig. 2. Two clear differences between samples are the thicknesses of the Er2O3 and the oxide layer of the F82H substrate. The sample withdrawn at 0.75 mm s1 showed approximately half the thickness of the Er2O3 layer (0.1 lm) and a thicker oxide layer of the substrate (up to 50 nm) in comparison with the sample withdrawn at 1.2 mm s1. In addition, the surface profile of the sample withdrawn at 0.75 mm s1 was rather rough because of the inhomogeneous oxide layer. Fig. 3 shows grazing-incidence XRD spectra of the coatings withdrawn at 0.75 (Sample 1), 1.2 mm s1 (Sample 2), and double processed at 1.2 mm s1 (twice that of the Sample 2 fabrication) (Sample 3). The reference peaks of the cubic Er2O3 are also presented [15]. Topmost surface information can be mainly detected by an incident angle of 0.5° corresponding to a submicron X-ray penetration depth. The sample withdrawn at 0.75 mm s1 showed the cubic Er2O3 peaks as well as the FeCr2O4 (3 1 1) peak at 35.5° derived from the oxide layer of the substrate [16]. On the other hand, the FeCr2O4 peak was fairly small on the sample withdrawn at 1.2 mm s1. In addition, the a-Fe (1 1 0) peak at 44.7° [17] from the F82H substrate was not detected on the double-coated sample, indicating an increase of Er2O3 thickness. Deuterium permeation experiments were performed for three representative samples. Arrhenius plots of the permeabilities of the coatings withdrawn at 0.75 and 1.2 mm s1 are shown in Fig. 4. The data for the F82H substrate is also provided [7]. Sample 1 showed the PRF of 20–30 at 773–973 K. On the other hand, Sample 2 showed a PRF of 600–700 at below 873 K, though the permeability increased by nearly one order of magnitude after the measurement at 923 K. Regarding Sample 3, the permeation experiment was commenced at 823 K for 1 day so that crystallization of the coating was complete. The temperature dependence of the permeabilities of Sample 3 is presented in Fig. 5. The first five plots at 773–973 K confirmed that the temperature dependence showed a linear relationship. It is estimated that the activation energy of permeation is 105 kJ mol1, which is approximately twice as much as that of the coating fabricated on one side of the substrate by spin coating [12].
Mirror-polished RAFM steel F82H (8Cr–2 W, heat No. 9753 42W-4) plates with dimensions of 25-mm length and 0.5-mm thickness were used as substrates. The MOD coating procedure is the same as in the previous study except for the coating method [12]. First, the substrate is dipped into an Er2O3 coating precursor (Kojundo Chemical Laboratory Co., Ltd. Er-03Ò) without addition of thinner, and then withdrawn at a constant speed of 0.05– 1.5 mm s1 by a dip-coater. Second, the sample is placed in a dry oven set at 393 K for 10 min to turn the solvent into a gel. The process of dipping and drying is repeated in order to reduce gravity-induced unevenness of the coating. Finally, the sample is heat treated in an infrared image furnace to crystallize the Er2O3 coating. In this study, the heat-treatment condition is determinedfrom the previous study using the spin-coating technique [12]: at 973 K for 10 min in hydrogen (purity: 99.9999%, O2: <0.02 ppm, H2O: <0.5 ppm) with approximately 0.6% moisture. The flow rate is less than 10 ml min1, and the rate of temperature increase and decrease is 100 and 30 K min1, respectively. Moisture is added to the hydrogen flow by passing it through ice-chilled water.
2.2. Coating characterization Surface profiles of the coated samples were examined with a confocal laser scanning microscope (CLSM). Cross sections of the samples prepared by a cross-section polisher were observed with a field-emission scanning electron microscope (FE-SEM). Crystal structures of the coatings were analyzed by X-ray diffraction (XRD). The configuration of instruments and the procedure for deuterium permeation experiments are described in detail in Refs. [7,10]. The sample is heated up to 773–973 K during the measurements so that a clear deuterium permeation can be detected by a quadrupole mass spectrometer (QMS). Driving pressures of deuterium are set to 1.00–8.00 104 Pa. The permeation flux per unit area at steady-state J (mol m2 s1) through a sample with a thickness of d (m) is expressed by the following equation:
J¼P
p0:5 ; d
ð1Þ
where P (mol m1 s1 Pa0.5) is the permeability, and p (Pa) is the driving pressure. An exponent of driving pressure is estimated by fitting the permeation fluxes at different driving pressures. The simplest case of permeation, where the rate-limiting process is diffusion of hydrogen atoms through the solid, namely a diffusion-limited regime (DLR), satisfies Eq. (1). On the other hand, when the permeation rate is limited by surface processes such as absorption and desorption, called a surface-limited regime (SLR), the exponent of the driving pressure p is unity. Therefore, in this study, if a deuterium pressure dependence of the permeation data is not applicable to the DLR, the permeability will be calculated with the driving pressure of 8.00 104 Pa, which is the highest pressure in this work, and then has a relatively small effect on surface reactions. In addition, as described in Ref. [10], the activation energy of the permeation can be evaluated by fitting a gradient of temperature dependence of the permeability because the permeation is a thermally-activated process. For evaluation of TPB efficiency, a permeation reduction factor (PRF) is calculated dividing the permeation flux of a bare substrate by that of a coated one. Reproducibility of the permeation data is confirmed by measuring several samples with the same coating parameter.
4. Discussion The relationship among fabrication parameters, coating properties, and deuterium permeation behaviors is discussed in this section. The withdrawal speed during the dipping process generally determines coating thickness; the higher the withdrawal speed, the thicker is the coating. In the case of the samples withdrawn at below 0.75 mm s1, the coatings might be too thin to be crystallized without oxidation of the F82H substrate during the heattreatment process. On the other hand, the coating was too thick to remain adhered to the substrate during the heat-treatment when withdrawn at 1.5 mm s1. The cross-sectional FE-SEM micrographs indicate that a proper coating thickness is essential to control formation of the oxide layer of the substrate, and the uniformity of the Er2O3 coating is strongly affected by the formation of the oxide layer. These arguments on the coating structure
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T. Chikada et al. / Journal of Nuclear Materials 442 (2013) 533–537
Fig. 1. Surface CLSM images of the coatings with withdrawal speeds of 0.05, 0.1, 0.5, 0.75, 1.0, 1.2, 1.4, and 1.5 mm s1.
0.75 mm s–1
1.2 mm s–1
Er2O3 Er2O3 FeCr2O4
FeCr2O4
F82H
F82H
100 nm
Fig. 2. Cross-sectional SEM images of the coatings withdrawn at 0.75 and 1.2 mm s1.
Cubic Er2O3
Intensity (a.u.)
Sample 3
Sample 2
FeCr2O4
Fe
Sample 1 Fig. 4. Arrhenius plots of deuterium permeability for Samples 1 and 2, and uncoated F82H [7].
20
30
40
50
60
Diffraction angle 2θ (º) Fig. 3. Grazing-incidence XRD spectra of the coatings with incident angle of 0.5°: withdrawn at 0.75 mm s1 (Sample 1), 1.2 mm s1 (Sample 2), and doubleprocessed at 1.2 mm s1 (twice that of the Sample 2 fabrication) (Sample 3).
are consistent with the previous study on the spin-coating technique [12]. The other feature of the coatings is the grain size, which was estimated as being approximately 4 nm from the Er2O3 peaks of the XRD spectra and Scherrer’s equation [18]. Since
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Fig. 6. Driving pressure dependence of deuterium permeation flux of Sample 3. Numbers in parentheses represent the pressure exponents calculated by fitting curves. Fig. 5. Arrhenius plots of deuterium permeability for Sample 3 and uncoated F82H [7].
an average grain size of 280 nm was confirmed in high-purity Er2O3 coatings fabricated by filtered arc deposition after the deuterium permeation experiment at 973 K [9], the formation of nanocrystal Er2O3 might be attributed to residual impurities such as carbon and other carbon compounds. The MOD coatings in this study did not show any difference in XRD pattern before and after permeation experiments up to 973 K. Deuterium permeation through the coated samples can be considered a sensitive sensor for degradation of the coatings. Comparing deuterium permeabilities of the samples with different withdrawal speeds (Fig. 4) clearly indicates that the uneven structure of the coating including the oxide layer of the substrate, easily generated pores and cracks that could be paths for permeation. The transition of the permeability seen in Sample 2 is supposed to include two phenomena: crystallization and crack formation. On the first measurements at 773 K, the permeation flux was continuously decreasing, indicating the Er2O3 coating was not completely crystallized. So the crystallization might be completed during the measurement at 823 K. This mechanism caused the different temperature dependence of the permeability between 773–823 and 823–873 K. The gradient of the temperature dependence from the second (823 K) to the third (873 K) permeation data does not indicate crack formation because the permeability at 823 K was reproduced after the measurement at 873 K. Therefore, a lower permeability is expected in the first data at 773 K if the measurement is conducted after crystallization at above 823 K. On the other hand, the steeper gradient between 873 and 923 K is due to the crack formation in the coating because the permeability was not reproduced after the measurement. The limiting process of deuterium permeation in the coatings could not be evaluated for these samples because the driving pressure dependence of the permeation flux was unclear. The difference in the activation energy of permeation between samples coated on one side and both sides was discussed in our previous studies [10,19]. According to those papers, when there are plural layers affecting the rate-limiting process of permeation, the activation energy of the sample is the sum of those of the layers. The activation energy of permeation obtained in this work ensures consistency with the discussion and also proves the integrity of the coating. However, after the measurement at 973 K, the permeability increased by one order of magnitude because of crack formation. Therefore, thermal durability of the coating remains an issue in the development of the dip-coating technique, although practical blanket systems with RAFM steels will not be applied at 973 K. Optimization of the number of layers and clarification of the permeation behavior under thermal cycles are required for further investigation. Besides, developing the dip-coating technique
for tubular substrates and evaluating their tritium permeability will be the next step to extrapolate the results in this study to actual utilization. A clear driving pressure dependence of deuterium permeation was obtained on Sample 3, as shown in Fig. 6. The pressure exponent tended to approach 0.5 at higher temperatures, indicating the transition of the rate-limiting process. Since a large gradient of temperature dependence of the permeability produces high PRFs at lower temperatures, surface reactions may strongly contribute to the suppression of deuterium permeation at low temperatures. On the other hand, tritium inventory in the coatings should be regarded, particularly at lower temperature regions. This phenomenon is a key difference from the both-sides-coated samples fabricated by filtered vacuum arc deposition [10]. Therefore the surface reactions may be induced by the oxide layer of the substrate and by impurities in the Er2O3 coating. Data accumulation and further consideration of not only hydrogen permeability but also diffusivity and solubility are required for a deeper understanding of tritium behavior in the coatings. 5. Summary Er2O3 coatings have been fabricated by dip coating RAFM steel F82H substrates. The coating parameter was optimized by controlling the withdrawal speed and the surface analysis. The 0.2-lmthick smooth coatings showed PRFs of 600–700 at 773 K; however, the PRF decreased by one order of magnitude because of crack formation after the permeation measurement at 923 K. The doublecoated sample showed a PRF of 2000 at 773 K and did not degrade at up to 923 K. Optimization of the number of repetitions of the coating procedure will be effective in reducing degradation and increasing thermal durability of the coatings. The clear relationship between permeability and temperature proved the reasonable value of activation energy of permeation. The strong contribution of surface reactions at low temperatures suggests the possibility of tritium accumulation in the coatings. Acknowledgement This work was supported in part by KAKENHI (23860017), Ministry of Education, Culture, Sports, Science and Technology, Japan. References [1] G.W. Hollenberg, E.P. Simonen, G. Kalinin, A. Terlain, Fusion Eng. Des. 28 (1995) 190–208. [2] A. Perujo, K.S. Forcey, Fusion Eng. Des. 28 (1995) 252–257. [3] G. Benamati, C. Chabrol, A. Perujo, E. Rigal, H. Glasbrenner, J. Nucl. Mater. 271 (272) (1999) 391–395.
T. Chikada et al. / Journal of Nuclear Materials 442 (2013) 533–537 [4] J. Konys, A. Aiello, G. Benamati, L. Giancarli, Fusion Sci. Technol. 47 (2005) 844– 850. [5] F. Koch, R. Brill, H. Maier, D. Levchuk, A. Suzuki, T. Muroga, H. Bolt, J. Nucl. Mater. 329–333 (2004) 1403–1406. [6] D. Levchuk, S. Levchuk, H. Maier, J. Nucl. Mater. 367–370 (2007) 1033–1037. [7] T. Chikada, A. Suzuki, Z. Yao, D. Levchuk, H. Maier, T. Terai, T. Muroga, Fusion Eng. Des. 84 (2009) 590–592. [8] T. Chikada, A. Suzuki, T. Kobayashi, Z. Yao, D. Levchuk, H. Maier, T. Terai, T. Muroga, Fusion Sci. Technol. 56 (2009) 309–313. [9] T. Chikada, A. Suzuki, T. Kobayashi, H. Maier, T. Terai, T. Muroga, J. Nucl. Mater. 417 (2011) 1241–1244. [10] T. Chikada, A. Suzuki, C. Adelhelm, T. Terai, T. Muroga, Nucl. Fusion 51 (2011) 063023.
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[11] Y. Hishinuma, T. Tanaka, T. Tanaka, T. Nagasaka, Y. Tasaki, S. Murakami, K. Matsuda, A. Sagara, T. Muroga, Fusion Sci. Technol. 60 (2011) 1131–1134. [12] T. Chikada, A. Suzuki, T. Tanaka, T. Terai, T. Muroga, Fusion Eng. Des. 85 (2010) 1537–1541. [13] D. Zhang, T. Tanaka, T. Muroga, J. Nucl. Mater. 417 (2011) 1249–1252. [14] D. Zhang, T. Tanaka, T. Muroga, Fusion Sci. Technol. 60 (2011) 1576–1579. [15] ICDD powder diffraction file No. 43-1007. [16] ICDD powder diffraction file No. 34-140. [17] ICDD powder diffraction file No. 6-696. [18] A. Patterson, Phys. Rev. 56 (1939) 978–982. [19] T. Chikada, A. Suzuki, H. Maier, T. Terai, T. Muroga, Fusion Sci. Technol. 60 (2011) 389–393.
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Journal of Nuclear Materials journal homepage: www.elsevier.com/locate/jnucmat
Deuterium permeation through erbium oxide coatings on RAFM steels by a dip-coating technique Takumi Chikada a,⇑, Shunya Naitoh b, Akihiro Suzuki c, Takayuki Terai a, Teruya Tanaka d, Takeo Muroga d a
Institute of Engineering Innovation, School of Engineering, The University of Tokyo, 2-11-16 Yayoi, Bunkyo-ku, Tokyo 113-8656, Japan Department of Systems Innovation, Faculty of Engineering, The University of Tokyo, 7-3-1 Hongo, Bunkyo-ku, Tokyo 113-8656, Japan c Nuclear Professional School, School of Engineering, The University of Tokyo, 2-22 Shirakata-shirane, Tokai, Naka, Ibaraki 319-1188, Japan d Fusion Systems Research Division, Department of Helical Plasma Research, National Institute for Fusion Science, 322-6 Oroshi, Toki, Gifu 509-5292, Japan b
a r t i c l e
i n f o
Article history: Available online 7 June 2013
a b s t r a c t A tritium permeation barrier is a promising solution for the problems of tritium loss and radiological safety in fusion blanket systems. In recent years, erbium oxide coatings have shown remarkable permeation reduction factors. One of the remaining issues for the coatings is the establishment of plant-scale fabrication. In this study, erbium oxide thin films have been fabricated by a dip-coating technique, which has the potential to coat a complex-shaped substrate, and deuterium permeation behavior in the coatings has been examined. Crack-free coatings were formed on a reduced activation ferritic/martensitic steel F82H substrate by use of a withdrawal speed of 1.0–1.4 mm s1 and a heat-treatment process in hydrogen with moisture. In deuterium permeation experiments, a 0.2-lm-thick coating on both sides of the substrate showed a reduction factor of 600–700 in comparison with a F82H substrate below 873 K; however, the coating degraded at above 923 K because of crack formation. A double-coated sample indicated a reduction factor of up to 2000 and did not degrade at up to 923 K. The driving pressure dependence of the deuterium permeation flux indicated that the permeation tended to be limited by surface reactions at low temperatures. Optimization of the number of layers has the possibility to reduce degradation at high temperatures while maintaining high permeation reduction factors. Ó 2013 Elsevier B.V. All rights reserved.
1. Introduction Control of tritium inventory and tritium leakage has gained attention for the development of deuterium–tritium (D–T) fusion reactors. One of the most crucial issues is tritium permeation through structural materials in blanket systems because of fatal fuel loss and radiological hazards. Ceramic coatings have been investigated for several decades to provide a thin hydrogenimpermeable layer on inner pipe walls as a tritium permeation barrier (TPB) [1–4]. In recent years, erbium oxide (Er2O3) coatings have been assessed as being a promising candidate for a TPB [5–7]. A series of studies achieved precise deuterium permeation behaviors through high-purity Er2O3 coatings deposited on reduced activation ferritic/martensitic (RAFM) steel substrates as well as permeation reduction factors (PRFs) of up to 105 at 873 K [8–10]. Recently, the status of the investigation has progressed to the stage of developing a fabrication process toward practical utilization such as plant-scale fabrication without restriction of s-
⇑ Corresponding author. Tel./fax: +81 3 5841 7420. E-mail address: [email protected] (T. Chikada). 0022-3115/$ - see front matter Ó 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jnucmat.2013.05.072
ubstrate geometry. One of the ongoing fabrication methods is metal–organic chemical vapor deposition (MOCVD). Fabrication of the Er2O3 coating on the tubular substrates was demonstrated. However, at present, a 1-lm-thick coating decreased hydrogen permeation to 1/20 of that with a SS316 plate [11]. Another is metal–organic decomposition (MOD), which is a liquid phase method with an Er-contained precursor. The liquid phase method has a great advantage in terms of simplicity of fabrication facilities (e.g., a large vacuum system is unnecessary), but control of coating properties such as crystallinity, the amount of impurities, and intermediate layers is challenging. In our previous study, a 0.3-lm-thick MOD sample coated by a spin-coating technique showed a PFR of 500–700 by controlling oxygen potential during the heat treatment [12]. However, in the case of the MOD samples coated by dip coating, approximately 1-lm-thick coatings on both sides of the RAFM substrate showed PRFs of up to 102 [13,14]. This result is probably due to the formation of cracks or peelings reaching the substrate even if the Er2O3 layer has been sufficiently crystallized. The goal of this study is to optimize the dip-coating parameter for the fabrication of crack-free MOD coatings and to acquire fine hydrogen permeation behaviors in the coatings.
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2. Experimental
3. Results
2.1. Sample preparation
As the first step in coating process optimization, surface profiles of the coated samples withdrawn at 0.05–1.5 mm s1 were compared, as shown in Fig. 1. The samples withdrawn at below 0.75 mm s1 showed spots. The sample withdrawn at 1.5 mm s1 showed obvious unevenness of the coating. No clear peeling was observed in the samples withdrawn at 1.0–1.4 mm s1. In particular smooth surfaces were observed on the samples withdrawn at 1.0–1.2 mm s1. Interfacial textures of the coatings were examined by cross-sectional observation, as shown in Fig. 2. Two clear differences between samples are the thicknesses of the Er2O3 and the oxide layer of the F82H substrate. The sample withdrawn at 0.75 mm s1 showed approximately half the thickness of the Er2O3 layer (0.1 lm) and a thicker oxide layer of the substrate (up to 50 nm) in comparison with the sample withdrawn at 1.2 mm s1. In addition, the surface profile of the sample withdrawn at 0.75 mm s1 was rather rough because of the inhomogeneous oxide layer. Fig. 3 shows grazing-incidence XRD spectra of the coatings withdrawn at 0.75 (Sample 1), 1.2 mm s1 (Sample 2), and double processed at 1.2 mm s1 (twice that of the Sample 2 fabrication) (Sample 3). The reference peaks of the cubic Er2O3 are also presented [15]. Topmost surface information can be mainly detected by an incident angle of 0.5° corresponding to a submicron X-ray penetration depth. The sample withdrawn at 0.75 mm s1 showed the cubic Er2O3 peaks as well as the FeCr2O4 (3 1 1) peak at 35.5° derived from the oxide layer of the substrate [16]. On the other hand, the FeCr2O4 peak was fairly small on the sample withdrawn at 1.2 mm s1. In addition, the a-Fe (1 1 0) peak at 44.7° [17] from the F82H substrate was not detected on the double-coated sample, indicating an increase of Er2O3 thickness. Deuterium permeation experiments were performed for three representative samples. Arrhenius plots of the permeabilities of the coatings withdrawn at 0.75 and 1.2 mm s1 are shown in Fig. 4. The data for the F82H substrate is also provided [7]. Sample 1 showed the PRF of 20–30 at 773–973 K. On the other hand, Sample 2 showed a PRF of 600–700 at below 873 K, though the permeability increased by nearly one order of magnitude after the measurement at 923 K. Regarding Sample 3, the permeation experiment was commenced at 823 K for 1 day so that crystallization of the coating was complete. The temperature dependence of the permeabilities of Sample 3 is presented in Fig. 5. The first five plots at 773–973 K confirmed that the temperature dependence showed a linear relationship. It is estimated that the activation energy of permeation is 105 kJ mol1, which is approximately twice as much as that of the coating fabricated on one side of the substrate by spin coating [12].
Mirror-polished RAFM steel F82H (8Cr–2 W, heat No. 9753 42W-4) plates with dimensions of 25-mm length and 0.5-mm thickness were used as substrates. The MOD coating procedure is the same as in the previous study except for the coating method [12]. First, the substrate is dipped into an Er2O3 coating precursor (Kojundo Chemical Laboratory Co., Ltd. Er-03Ò) without addition of thinner, and then withdrawn at a constant speed of 0.05– 1.5 mm s1 by a dip-coater. Second, the sample is placed in a dry oven set at 393 K for 10 min to turn the solvent into a gel. The process of dipping and drying is repeated in order to reduce gravity-induced unevenness of the coating. Finally, the sample is heat treated in an infrared image furnace to crystallize the Er2O3 coating. In this study, the heat-treatment condition is determinedfrom the previous study using the spin-coating technique [12]: at 973 K for 10 min in hydrogen (purity: 99.9999%, O2: <0.02 ppm, H2O: <0.5 ppm) with approximately 0.6% moisture. The flow rate is less than 10 ml min1, and the rate of temperature increase and decrease is 100 and 30 K min1, respectively. Moisture is added to the hydrogen flow by passing it through ice-chilled water.
2.2. Coating characterization Surface profiles of the coated samples were examined with a confocal laser scanning microscope (CLSM). Cross sections of the samples prepared by a cross-section polisher were observed with a field-emission scanning electron microscope (FE-SEM). Crystal structures of the coatings were analyzed by X-ray diffraction (XRD). The configuration of instruments and the procedure for deuterium permeation experiments are described in detail in Refs. [7,10]. The sample is heated up to 773–973 K during the measurements so that a clear deuterium permeation can be detected by a quadrupole mass spectrometer (QMS). Driving pressures of deuterium are set to 1.00–8.00 104 Pa. The permeation flux per unit area at steady-state J (mol m2 s1) through a sample with a thickness of d (m) is expressed by the following equation:
J¼P
p0:5 ; d
ð1Þ
where P (mol m1 s1 Pa0.5) is the permeability, and p (Pa) is the driving pressure. An exponent of driving pressure is estimated by fitting the permeation fluxes at different driving pressures. The simplest case of permeation, where the rate-limiting process is diffusion of hydrogen atoms through the solid, namely a diffusion-limited regime (DLR), satisfies Eq. (1). On the other hand, when the permeation rate is limited by surface processes such as absorption and desorption, called a surface-limited regime (SLR), the exponent of the driving pressure p is unity. Therefore, in this study, if a deuterium pressure dependence of the permeation data is not applicable to the DLR, the permeability will be calculated with the driving pressure of 8.00 104 Pa, which is the highest pressure in this work, and then has a relatively small effect on surface reactions. In addition, as described in Ref. [10], the activation energy of the permeation can be evaluated by fitting a gradient of temperature dependence of the permeability because the permeation is a thermally-activated process. For evaluation of TPB efficiency, a permeation reduction factor (PRF) is calculated dividing the permeation flux of a bare substrate by that of a coated one. Reproducibility of the permeation data is confirmed by measuring several samples with the same coating parameter.
4. Discussion The relationship among fabrication parameters, coating properties, and deuterium permeation behaviors is discussed in this section. The withdrawal speed during the dipping process generally determines coating thickness; the higher the withdrawal speed, the thicker is the coating. In the case of the samples withdrawn at below 0.75 mm s1, the coatings might be too thin to be crystallized without oxidation of the F82H substrate during the heattreatment process. On the other hand, the coating was too thick to remain adhered to the substrate during the heat-treatment when withdrawn at 1.5 mm s1. The cross-sectional FE-SEM micrographs indicate that a proper coating thickness is essential to control formation of the oxide layer of the substrate, and the uniformity of the Er2O3 coating is strongly affected by the formation of the oxide layer. These arguments on the coating structure
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Fig. 1. Surface CLSM images of the coatings with withdrawal speeds of 0.05, 0.1, 0.5, 0.75, 1.0, 1.2, 1.4, and 1.5 mm s1.
0.75 mm s–1
1.2 mm s–1
Er2O3 Er2O3 FeCr2O4
FeCr2O4
F82H
F82H
100 nm
Fig. 2. Cross-sectional SEM images of the coatings withdrawn at 0.75 and 1.2 mm s1.
Cubic Er2O3
Intensity (a.u.)
Sample 3
Sample 2
FeCr2O4
Fe
Sample 1 Fig. 4. Arrhenius plots of deuterium permeability for Samples 1 and 2, and uncoated F82H [7].
20
30
40
50
60
Diffraction angle 2θ (º) Fig. 3. Grazing-incidence XRD spectra of the coatings with incident angle of 0.5°: withdrawn at 0.75 mm s1 (Sample 1), 1.2 mm s1 (Sample 2), and doubleprocessed at 1.2 mm s1 (twice that of the Sample 2 fabrication) (Sample 3).
are consistent with the previous study on the spin-coating technique [12]. The other feature of the coatings is the grain size, which was estimated as being approximately 4 nm from the Er2O3 peaks of the XRD spectra and Scherrer’s equation [18]. Since
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Fig. 6. Driving pressure dependence of deuterium permeation flux of Sample 3. Numbers in parentheses represent the pressure exponents calculated by fitting curves. Fig. 5. Arrhenius plots of deuterium permeability for Sample 3 and uncoated F82H [7].
an average grain size of 280 nm was confirmed in high-purity Er2O3 coatings fabricated by filtered arc deposition after the deuterium permeation experiment at 973 K [9], the formation of nanocrystal Er2O3 might be attributed to residual impurities such as carbon and other carbon compounds. The MOD coatings in this study did not show any difference in XRD pattern before and after permeation experiments up to 973 K. Deuterium permeation through the coated samples can be considered a sensitive sensor for degradation of the coatings. Comparing deuterium permeabilities of the samples with different withdrawal speeds (Fig. 4) clearly indicates that the uneven structure of the coating including the oxide layer of the substrate, easily generated pores and cracks that could be paths for permeation. The transition of the permeability seen in Sample 2 is supposed to include two phenomena: crystallization and crack formation. On the first measurements at 773 K, the permeation flux was continuously decreasing, indicating the Er2O3 coating was not completely crystallized. So the crystallization might be completed during the measurement at 823 K. This mechanism caused the different temperature dependence of the permeability between 773–823 and 823–873 K. The gradient of the temperature dependence from the second (823 K) to the third (873 K) permeation data does not indicate crack formation because the permeability at 823 K was reproduced after the measurement at 873 K. Therefore, a lower permeability is expected in the first data at 773 K if the measurement is conducted after crystallization at above 823 K. On the other hand, the steeper gradient between 873 and 923 K is due to the crack formation in the coating because the permeability was not reproduced after the measurement. The limiting process of deuterium permeation in the coatings could not be evaluated for these samples because the driving pressure dependence of the permeation flux was unclear. The difference in the activation energy of permeation between samples coated on one side and both sides was discussed in our previous studies [10,19]. According to those papers, when there are plural layers affecting the rate-limiting process of permeation, the activation energy of the sample is the sum of those of the layers. The activation energy of permeation obtained in this work ensures consistency with the discussion and also proves the integrity of the coating. However, after the measurement at 973 K, the permeability increased by one order of magnitude because of crack formation. Therefore, thermal durability of the coating remains an issue in the development of the dip-coating technique, although practical blanket systems with RAFM steels will not be applied at 973 K. Optimization of the number of layers and clarification of the permeation behavior under thermal cycles are required for further investigation. Besides, developing the dip-coating technique
for tubular substrates and evaluating their tritium permeability will be the next step to extrapolate the results in this study to actual utilization. A clear driving pressure dependence of deuterium permeation was obtained on Sample 3, as shown in Fig. 6. The pressure exponent tended to approach 0.5 at higher temperatures, indicating the transition of the rate-limiting process. Since a large gradient of temperature dependence of the permeability produces high PRFs at lower temperatures, surface reactions may strongly contribute to the suppression of deuterium permeation at low temperatures. On the other hand, tritium inventory in the coatings should be regarded, particularly at lower temperature regions. This phenomenon is a key difference from the both-sides-coated samples fabricated by filtered vacuum arc deposition [10]. Therefore the surface reactions may be induced by the oxide layer of the substrate and by impurities in the Er2O3 coating. Data accumulation and further consideration of not only hydrogen permeability but also diffusivity and solubility are required for a deeper understanding of tritium behavior in the coatings. 5. Summary Er2O3 coatings have been fabricated by dip coating RAFM steel F82H substrates. The coating parameter was optimized by controlling the withdrawal speed and the surface analysis. The 0.2-lmthick smooth coatings showed PRFs of 600–700 at 773 K; however, the PRF decreased by one order of magnitude because of crack formation after the permeation measurement at 923 K. The doublecoated sample showed a PRF of 2000 at 773 K and did not degrade at up to 923 K. Optimization of the number of repetitions of the coating procedure will be effective in reducing degradation and increasing thermal durability of the coatings. The clear relationship between permeability and temperature proved the reasonable value of activation energy of permeation. The strong contribution of surface reactions at low temperatures suggests the possibility of tritium accumulation in the coatings. Acknowledgement This work was supported in part by KAKENHI (23860017), Ministry of Education, Culture, Sports, Science and Technology, Japan. References [1] G.W. Hollenberg, E.P. Simonen, G. Kalinin, A. Terlain, Fusion Eng. Des. 28 (1995) 190–208. [2] A. Perujo, K.S. Forcey, Fusion Eng. Des. 28 (1995) 252–257. [3] G. Benamati, C. Chabrol, A. Perujo, E. Rigal, H. Glasbrenner, J. Nucl. Mater. 271 (272) (1999) 391–395.
T. Chikada et al. / Journal of Nuclear Materials 442 (2013) 533–537 [4] J. Konys, A. Aiello, G. Benamati, L. Giancarli, Fusion Sci. Technol. 47 (2005) 844– 850. [5] F. Koch, R. Brill, H. Maier, D. Levchuk, A. Suzuki, T. Muroga, H. Bolt, J. Nucl. Mater. 329–333 (2004) 1403–1406. [6] D. Levchuk, S. Levchuk, H. Maier, J. Nucl. Mater. 367–370 (2007) 1033–1037. [7] T. Chikada, A. Suzuki, Z. Yao, D. Levchuk, H. Maier, T. Terai, T. Muroga, Fusion Eng. Des. 84 (2009) 590–592. [8] T. Chikada, A. Suzuki, T. Kobayashi, Z. Yao, D. Levchuk, H. Maier, T. Terai, T. Muroga, Fusion Sci. Technol. 56 (2009) 309–313. [9] T. Chikada, A. Suzuki, T. Kobayashi, H. Maier, T. Terai, T. Muroga, J. Nucl. Mater. 417 (2011) 1241–1244. [10] T. Chikada, A. Suzuki, C. Adelhelm, T. Terai, T. Muroga, Nucl. Fusion 51 (2011) 063023.
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[11] Y. Hishinuma, T. Tanaka, T. Tanaka, T. Nagasaka, Y. Tasaki, S. Murakami, K. Matsuda, A. Sagara, T. Muroga, Fusion Sci. Technol. 60 (2011) 1131–1134. [12] T. Chikada, A. Suzuki, T. Tanaka, T. Terai, T. Muroga, Fusion Eng. Des. 85 (2010) 1537–1541. [13] D. Zhang, T. Tanaka, T. Muroga, J. Nucl. Mater. 417 (2011) 1249–1252. [14] D. Zhang, T. Tanaka, T. Muroga, Fusion Sci. Technol. 60 (2011) 1576–1579. [15] ICDD powder diffraction file No. 43-1007. [16] ICDD powder diffraction file No. 34-140. [17] ICDD powder diffraction file No. 6-696. [18] A. Patterson, Phys. Rev. 56 (1939) 978–982. [19] T. Chikada, A. Suzuki, H. Maier, T. Terai, T. Muroga, Fusion Sci. Technol. 60 (2011) 389–393.