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Morphology evolution of La2O3 and crack characteristic in W–La2O3 alloy under transient heat loading
Journal of Nuclear Materials 451 (2014) 283–291
Contents lists available at ScienceDirect
Journal of Nuclear Materials journal homepage: www.elsevier.com/locate/jnucmat
Morphology evolution of La2O3 and crack characteristic in W–La2O3 alloy under transient heat loading Xiaoxin Zhang, Qingzhi Yan ⇑ Institute of Nuclear Materials, University of Science & Technology Beijing, 30 Xueyuan Road, Haidian District, Beijing, China
h i g h l i g h t s 2
Thermal shock tests were performed on un-heated WL10 samples at 0.22–1.1 GW/m . 2
The surfaces loaded at 0.88 and 1.1 GW/m were divided into four areas. In the transition area, the left vacancies aggravated the crack initiation. In the resolidification area, the craters did not facilitate the crack initiation. WL10 showed worse performance than pure tungsten when exposed to thermal loads.
a r t i c l e
i n f o
Article history: Received 21 October 2013 Accepted 1 April 2014 Available online 13 April 2014
a b s t r a c t Thermal shock tests were performed on W–La2O3 specimens by an electron beam to investigate the morphology evolution of La2O3 and crack characteristics. Cracks appeared at power densities ranging from 0.22 to 0.88 GW/m2 and disappeared completely at 1.1 GW/m2. The specimens tested at 0.22–0.66 GW/m2 kept their original morphologies except for cracks. For the specimens loaded at 0.88 and 1.1 GW/m2, the sample surfaces were divided into four areas including no heat-loaded area, two transition areas and resolidification area. In the transition areas, the left vacancies after melting, evaporation and dissociation of La2O3 particles kept their original elongated shape and aggravated the formation of crack. In the resolidification area, the craters were partly or completely filled by tungsten and the round craters with smooth edges did not facilitate the crack initiation. Thermal fatigue tests suggested that WL10 showed worse performance than pure tungsten when exposed to thermal loads. Ó 2014 Elsevier B.V. All rights reserved.
1. Introduction Tungsten (W) is considered as one of the most promising candidate for plasma facing materials (PFMs) in future fusion reactors because of its high melting point, high thermal conductivity, low deuterium/tritium retention and low sputter rates [1–4]. PFM will be subjected to heavy thermal loads: quasi-stationary heat fluxes, edge localized modes (ELMs), vertical displacement events and plasma disruptions [5–8]. W–La2O3 have been proved to exhibit better mechanical properties, higher recrystallization temperature, more toughness than pure tungsten [9–14]. So W–La2O3 is generally considered as structural material in the plasma facing component, while pure tungsten is considered as a plasma facing material, which will suffer high thermal loads [15]. To qualify the resistance to high thermal loads of tungsten, high heat load tests have been performed in plasma simulators, neutral beam facility ⇑ Corresponding author. Tel./fax: +86 01062334951. E-mail address: [email protected] (Q. Yan). http://dx.doi.org/10.1016/j.jnucmat.2014.04.001 0022-3115/Ó 2014 Elsevier B.V. All rights reserved.
and electron beam facility [16–22]. In Ref. [16], mock-ups armored with W rods (pure W and lanthanated W) embedded in watercooled copper-alloy heat sinks were prepared, and these mockups exhibited excellent performance in thermal response tests at up to 30 MW/m2 and in thermal cycling tests of 500 cycles (10 s on, 10 s off) at 25 MW/m2. In Ref. [17], the integral bonding of tungsten tube with an internal 38-mm-long tungsten porous mesh was demonstrated and this integrally bonded tungsten heat sink displayed outstanding heat removal with a peak absorbed heat flux of 22.4 MW/m2 with He at 4 MPa, flowing at 27 g/s. Especially, Refs. [18–22] showed the different performance of W–La2O3 under high heat flux tests. In Ref. [18], W–La2O3 samples were irradiated by powerful deuterium plasma shots and the doped La2O3 particles decrease the development of a mesh of microcracks on the irradiated surface, which was conditioned by the greater brittleness of pure tungsten in comparison with W–La2O3; In Ref. [19], heat loads were applied to pure tungsten and W–La2O3 samples in the neutral beam facility GLADIS. Bubble formation was greatly suppressed in W–La2O3 than pure tungsten,
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X. Zhang, Q. Yan / Journal of Nuclear Materials 451 (2014) 283–291
Fig. 1. Microstructures of the rolled WL10 on cross-section surface (a–c) and the illustration of electron beam orientation with respect to the loaded surface and rolling direction (d), RD, TD and ND stand for rolling direction, transverse direction and normal direction, respectively.
Absorbed energy, J
10
8
6
4
2
600
700
Fig. 2. Surface temperatures of WL10 during thermal shock and fatigue tests.
800
900
1000
1100
1200
1300
Temperature, K Fig. 3. Results of the Charpy impact tests for WL10.
which was attributed to the reduced boiling of the molten tungsten by the lower melting threshold of W–La2O3 compared with pure tungsten and the consumption of the required latent heat of the evaporated La2O3. However, some opposite results reported for the thermal shock resistance of W–La2O3. In Ref. [20], behavior of W–La2O3 under plasma disruption simulation conditions was investigated in 2MK-200 facility and the wave surface structure was caused by the low melting temperature of the La2O3 particles; In Ref. [21], thermal shock tests were performed on W–La2O3 specimens with absorbed power density up to 1.1 GW/m2 for pulse
duration of 4.6 ms and significant release of LaO molecules, melting of the surfaces were clearly detected at 0.66 GW/m2 and 1.1 GW/m2, respectively. Besides, severe cracking was observed at the vicinity of the melting spots. So the authors indicated that adding La2O3 in tungsten did not mean the improvement of the thermal shock resistance and it caused introduction of additional impurity in vacuum chamber. However, the morphological evolution of La2O3 particles was not described in this study. Ghezzi
Table 1 Thermal–mechanical properties and detailed chemical impurities information of the annealed WL10 sample. Sample
Relative density (%)
WL10
97.2 452 1312 140.1 Impurities (ppm): C < 20, S < 20, Fe < 15, Ni < 10, Al < 15, Cr < 10, Cu < 10, Ca < 20, Mg < 10, Bi < 1, Sn < 5, Si < 10, K < 10, Pb < 1, As < 10
Micro-hardness (HV)
Bending strength (MPa)
Thermal conductivity (W/mK)
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Fig. 4. Photos of WL10 after the Charpy impact tests.
Fig. 5. An overview of the thermal shock response of the rolled WL10 exposed to the electron beam.
et al. examined the surface and bulk modification of W–La2O3 exposed to 100 ELM’s transient loads at the QSPA facility, but they more focused on the surface modification of tungsten as well as the evolution of La2O3 particle density from the bulk to the surface [22]. In summary, the morphological evolution of La2O3 particles under transient heat loading has not been reported and the effect of doping La2O3 on the thermal load resistance of tungsten is not completely clear by now. In present work, we focused on the morphological evolution of La2O3 particles as well as other thermal loading responses of tungsten in terms of crack characteristics and surface roughness.
2. Experimental The investigated W-1 wt% La2O3 (WL10) was prepared by powder metallurgical route followed by rolling process. The sintered billets were rolled from 25-mm-thick to 12 mm as the rolling temperature changed from the initial 1873 K to the target 1573 K with the engineering strain of 52%. The rolled samples were annealed at 1373 K in hydrogen atmosphere for 2 h to relax residual stress. Then the samples were cut into the size of 12 12 3 mm3 and polished to remove cutting induced scratches and other defects. Finally, the prepared specimens were mounted into a copper sample holder, with a heat removal system on the
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Fig. 6. SEM images of WL10 samples tested at 0.88 and 1.1 GW/m2.
Fig. 7. High magnification SEM images of slightly melted La2O3 in area B.
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Fig. 8. High magnification SEM images of the heavily melted La2O3 in area C.
reverse. The grain orientation was parallel to the heat transfer direction (see Fig. 1(d)), which corresponds to the ITER specification. Thermal shock tests were performed at absorbed power densities varied from 0.22 to 1.1 GW/m2 at an increment of 0.22 GW/ m2 with just 1 pulse. The detailed parameters for the transient high heat load tests were described in Ref. [23]. A homogeneous heat load distribution in the 4 4 mm2 beam spot was achieved by fast scanning (10 kHz) of the electron beam. Fig. 2 shows the surface temperature evolution of WL10 during thermal shock and fatigue tests. The surface temperature showed a general increase with the absorbed power density except 0.88 GW/m2 during the thermal shock tests, which should be caused by the melting of La2O3 particles. For the thermal fatigue tests, surface temperatures were 600–800 K. For Charpy impact tests, specimens with dimensions of 27 4 3 mm (length width height), a notch depth of 1 mm and a notch root radius of 0.1 mm were prepared. The notch orientation was R–N (‘‘R’’ stands for the rolling direction; ‘‘N’’ stands for the direction of the thickness of the plate. The first letter indicates the direction perpendicular to the expected crack plane while the
second letter stands for the expected direction of crack growth). Reiser also reported that this orientation resulted in optimum Charpy properties [24]. The specimens were heated to 673– 1273 K in argon atmosphere, then pushed to the support outside the furnace and hit by a striker (25 J) within about 2 s. The temperature drop was around 50 K measured by an infrared thermometer (HT-8865, China). Microstructure obversion and surface analysis were carried out with an optical microscope and scanning electron microscopy (SEM). Vickers micro-hardness tests were performed on the polished surface under a load of 1.96 N for 10 s by a MH-6 micro-hardness instrument. Three point bending (3PB) specimens of RD(37) TD(4) ND(3) were produced to measure the fracture strength. The values in the brackets stand for the dimensions along different directions. The 3PB tests were performed at room temperature with the deformation speed of 0.5 mm/min. Surface roughness was measured along five paths over the loaded surface by a laser profilometry. The thermal conductivity was calculated by the following formula:
k ¼ q cp a
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Fig. 9. SEM images of left vacancies in the resolidification area, low magnification image (a) and high magnification images (b–e) corresponding to the selected square region in Fig. 9 (a).
where k is the thermal conductivity, q is bulk density, cp is specific heat, a is the thermal diffusivity. The specific heat was measured by a thermal analyzing apparatus (Dupont 1090B, America). The thermal diffusivity was determined using a laser flash apparatus (LFA 457, Germany). Density was measured using Archimedes’ method.
3. Results and discussion Microstructures of the annealed WL10 are shown in Fig. 1. The pancake-like grains are flattened parallel to the plate surfaces (Fig. 1(a)), the needle-shaped grains are extremely elongated along the rolling direction with an aspect ratio of roughly 3.5:1 (Fig. 1(c)). Table 1 presents the thermal–mechanical properties and detailed chemical impurities information of the annealed WL10 sample. The thermal conductivity of 140.1 W/mK is higher than 121 W/
mK reported in Ref. [25], which should be attributed to the higher relative density of 97.2% compared with 95.0% in Ref. [25]. Fig. 3 shows the results of the Charpy impact tests and the photos of WL10 samples after the Charpy impact tests are presented in Fig. 4. WL10 specimens showed delamination (fractures propagate parallel to the specimen’s long side and perpendicular to the notch) at temperatures ranging from 773 K to 1273 K. The delamination of WL10 in Charpy test was also observed in Ref. [26]. According to the Charpy test results, the brittle-to-delamination transition temperature (defined in analogy to the ductile-to-brittle transition temperature (DBTT)) was 723–773 K for WL10. Thermal shock tests were carried out on un-heated WL10 samples at absorbed power densities varied from 0.22 to 1.1 GW/m2 in a step of 0.22 GW/m2 with grains elongated orientation paralleling to the heat flux. An overview of the thermal shock response of the annealed WL10 exposed to electron beam is given in Fig. 5. Cracks
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occurred in the samples loaded at power densities of 0.22– 0.88 GW/m2 and disappeared completely in the sample tested at the highest power density of 1.1 GW/m2. Furthermore, cracks displayed circular profile and branched at lower power density including 0.22, 0.44, 0.66 GW/m2. The reasons for the circular crack profile were described in Ref. [23]. But for the sample loaded at 0.88 GW/m2, the phenomenon of circular cracks branching was not detected. In other words, the crack density decreased as the power density increased. A similar phenomenon has been reported in Ref. [27], in which the microcrack networks were replaced by discontinuing networks and also disappear completely at higher power densities. Firstly, we investigated the amazing samples loaded at 0.88 GW/m2 and 1.1 GW/m2. Melting of tungsten matrix in samples loaded above 0.88 GW/m2 was detected obviously. Fig. 6 shows the SEM images of WL10 samples tested at 0.88 and 1.1 GW/m2. Basing on the distance from the center of scanning area, there are four different areas on the surface: no heat-loaded area (marked as A), two transition areas (marked as B and C), resolidification area (marked as D) with a diameter of 1–1.5 mm. In area A, the surface was not scanned by electron beam and kept the original morphology before thermal shock tests; In area B, C, the surface has been scanned homogeneously by electron beam but B area exhibited more brightness compared with C area; D represents the center of electron scanning area, the molten tungsten matrix solidified obviously. Fig. 6(c), (d), (e), (f) are high magnification images corresponding to A, B, C, D areas in Fig. 6(a) respectively. A area almostly kept the morphology before tests except cracks (Fig. 6(c)), which was caused by the temperature gradients and the related thermal stresses. In B area, La2O3 particles slightly melted and kept its original elongated shape (Figs. 6(d) and 7). Fig. 7 shows the high magnification SEM images of slightly melted La2O3 in area B, and tungsten element, oxygen element and lanthanum element were not detected resulting from the depth of left vacancies after the evaporation, dissociation of La2O3 particles went beyond the measurement range of SEM. The dissociation of La2O3 particles at 2773 K during thermal shock process has been proved in Ref. [21] and the evaporization of La2O3 particles must occured in samples loaded at 0.66, 0.88 and 1.1 GW/m2 with the surface temperatures over the melting point of La2O3 (2573 K). In area C, La2O3 particles melted heavily, then the melted La2O3 particles overflowed and covered the near tungsten matrix (Figs. 6(e) and 8). Fig. 8 shows the high magnification SEM images of the melted La2O3 in transition area corresponding to the selected square region in Fig. 6(e). The overflowed La2O3 particles were darker than tungsten matrix, which results in more brightness of area B compared with area C. In the resolidification area D, tungsten matrix melted obviously, then solidified and formed hill-like
structure (Figs. 6(f) and 9). Fig. 9 shows the high magnification SEM images of the resolidification area. Besides, the elongated left vacancies were replaced by the craters, which should be caused by the following process: (1) La2O3 particles as well as the left vacancies were covered with the molten tungsten, (2) gases originated from the evaporation, dissociation of La2O3 particles escaped into the ambient resulting in the craters formation, and (3) the craters were partly or completely filled with the molten tungsten (Fig. 9). More importantly, the left vacancies after melting, evaporation and dissociation of La2O3 particles influenced the crack formation significantly. The elongated left vacancies in the transition areas with tip aggravated crack formation as the stress concentration during the cooling phase (see Fig. 10). In the resolidification area, the craters with smooth edges did not facilitate the crack initiation. Secondly, the surface roughness and the crack characteristics in terms of crack depth and crack width were investigated. Fig. 11 shows the arithmetic mean roughness (Ra) as a function of power density. The surface roughness of loaded area increases slightly to a certain value from 0.22 to 0.44 GW/m2, stays constant (about 140 nm) from 0.44 to 0.66 GW/m2 and finally increase to a maximum value (about 320 nm) at 1.1 GW/m2. In the resolidification area, the hills developed obviously and the corresponding roughness values were about 6500 and 1300 nm for the samples loaded at 0.88 and 1.1 GW/m2 respectively. Besides, the maximal crack width and mean crack depth were quantified from the microstructures and depicted in Fig. 12. The maximal crack width (around
Fig. 10. SEM image of crack initiation in transition area.
Fig. 12. Power density dependence of crack parameters.
Fig. 11. Arithmetic mean roughness as a function of power density.
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Fig. 13. SEM images of the rolled WL10 after thermal fatigue tests.
8 lm) and depth (around 900 lm) were observed at 0.66 and 0.88 GW/m2, respectively. In the range from 0.22 to 0.88 GW/m2, both crack parameters showed a general increase with power density caused by the increase of the temperature gradients and the related thermal stresses as well as the elongated vacancies after melting, evaporation and dissociation of La2O3 particles. As the power density increases above 0.66 GW/m2, both crack parameters tended to minish and the cracks even disappeared at the sample loaded at 1.1 GW/m2, which should be caused by the decrease of the elastic modulus at the elevated temperature and the degradation of mechanical properties. Finally, thermal fatigue tests were performed on WL10 samples at an absorbed power density of 0.17 GW/m2 with pulse up to 1000. No deterioration was observed after 1 pulse which indicated the cracking threshold for WL10 was in the range of 0.17–0.22 GW/ m2. An overview of the thermal fatigue response of the rolled WL10 exposed to electron beam with 100, 1000 pulses is given in Fig. 13(a) and (b). After 100 pulses, the cracks with the maximal width of 3 lm appeared and the La2O3 particles kept their original shape without deterioration (Fig. 13(c)). After 1000 pulses, the maximal crack width increased up to 9 lm. Besides, the roughening of La2O3 particles was detected (Fig. 13(d)), which should be caused by the evaporation of La2O3 particles. In comparison with the responses of pure tungsten in Ref. [23], WL10 samples showed a worse performance than pure tungsten when exposed to thermal loads. The different responses of tungsten grades including pure tungsten and WL10 as well as potassium doped tungsten under transient heat loading will be explained in future. 4. Conclusions Thermal shock and fatigue tests were performed on specimens cut from a rolled WL10 in an electron beam facility and the grain orientation was parallel to the heat transfer direction. The brittle-to-delamination transition temperature was 723–773 K for WL10. Thermal shock tests were performed at absorbed power densities varied from 0.22 to 1.1 GW/m2 with an increment of 0.22 GW/m2. Cracks appeared in the samples loaded at power densities ranging from 0.22 to 0.88 GW/m2 and disappeared com-
pletely in the sample tested at the highest power density of 1.1 GW/m2. The specimens tested at 0.22–0.66 GW/m2 kept their original morphologies except for cracks. For the specimens loaded at 0.88 and 1.1 GW/m2, the sample surfaces were divided into four areas including no heat-loaded area, two transition areas and resolidification area. In the transition areas, the melted La2O3 overflowed, covered the near tungsten matrix and the elongated left vacancies aggravated the formation of crack as the stress concentration during the cooling phase. In the resolidification area, La2O3 particles and tungsten matrix suffered the greatest melt. The elongated left vacancies were replaced by the craters and the craters were partly or completely filled with the molten tungsten resulting in smooth edges, which did not facilitate the crack initiation. Finally, thermal fatigue tests were performed at absorbed power density of 0.17 GW/m2 with pulse up to 1000, the maximal crack widths were 3 lm and 9 lm when the pulses increased from 100 to 1000. All the findings suggested that WL10 showed a worse performance than pure tungsten when exposed to thermal loads. Acknowledgement The authors gratefully acknowledge the financial support of ITER-National Magnetic Confinement Fusion Program (2014GB123000). References [1] M. Rieth et al., J. Nucl. Mater. 417 (2011) 463–467. [2] V. Philipps, J. Nucl. Mater. 415 (2011) S2–S9. [3] T. Hao, Z.Q. Fan, S.X. Zhao, G.N. Luo, C.S. Liu, Q.F. Fang, J. Nucl. Mater. 433 (2012) 351–356. [4] Z. Zhou, J. Tan, D. Qu, G. Pintsuk, M. RöDig, J. Linke, J. Nucl. Mater. 431 (2011) 202–205. [5] M. Wirtz, J. Linke, G. Pintsuk, L. Singheiser, M. Zlobinski, J. Nucl. Mater. 438 (2013) S833–S836, http://dx.doi.org/10.1016/j.jnucmat.2013.01.180. [6] G. Pintsuk, A. Prokhodtseva, I. Uytdenhouwen, J. Nucl. Mater. 417 (2011) 481– 486. [7] M. Merola, M. RöDig, J. Linke, R. Duwe, G. Vieider, J. Nucl. Mater. 258 (1998) 653–657. [8] M. Roedig, I. Kupriyanov, J. Linke, X. Liu, Z. Wang, J. Nucl. Mater. 417 (2011) 761–764. [9] Q. Yan, X. Zhang, T. Wang, C. Yang, C. Ge, J. Nucl. Mater. 442 (2013) S233–S236, http://dx.doi.org/10.1016/j.jnucmat.2013.01.307.
X. Zhang, Q. Yan / Journal of Nuclear Materials 451 (2014) 283–291 [10] M. Faleschini, H. Kreuzer, D. Kiener, R. Pippan, J. Nucl. Mater. 367 (2007) 800– 805. [11] M. Mabuchi, K. Okamoto, N. Saito, M. Nakanishi, Y. Yamada, T. Asahina, T. Igarashi, Mater. Sci. Eng. A 214 (1996) 174–176. [12] M. Mabuchi, K. Okamoto, N. Saito, T. Asahina, T. Igarashi, Mater. Sci. Eng. A 237 (1997) 241–249. [13] I. Wesemann, W. Spielmann, P. Heel, A. Hoffmann, Int. J. Refract. Met. Hard Mater. 28 (2010) 687–691. [14] Y. Kim, K.H. Lee, E. Kim, D. Cheong, S.H. Hong, Int. J. Refract. Met. Hard Mater. 27 (2009) 842–846. [15] R.E. Nygren, R. Raffray, D. Whyte, M.A. Urickson, C.P.C. Wong, M. Baldwin, L.L. Snead, J. Nucl. Mater. 417 (2011) 451–456. [16] R.E. Nygren, D.L. Youchison, R.D. Watson, Eng. Des. 49–50 (2000) 303–308. [17] D.L. Youchison, T.J. Lutz, B. Williams, R.E. Nygren, Eng. Des. 82 (2007) 1854– 1860. [18] V. Shirokova, T. Laas, A. Ainsaar, J. Priimets, Ü. Ugaste, E.V. Demina, V.N. Pimenov, S.A. Maslyaev, A.V. Dubrovsky, V.A. Gribkov, J. Nucl. Mater. 435 (2012) 181–188.
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[19] Y. Yuan, H. Greuner, B. BöSwirth, G. Luo, B.Q. Fu, H.Y. Xu, W. Liu, J. Nucl. Mater. 438 (2013) S229–S232, http://dx.doi.org/10.1016/j.jnucmat.2013.01.033. [20] A. Makhankov, V. Barabash, I. Mazul, D. Youchison, J. Nucl. Mater. 290 (2001) 1117–1122. [21] T. Hirai, W. Kühnlein, J. Linke, G. Sergienko, Dunamic erosion of plasma facing materials under ITER relevant thermal shock loads in the electron beam facility, JUDITH, in: 20th IAEA Fusion Energy Conference, 2004-111-2004-11-6 Vilamoura, Portugal. [22] F. Ghezzi, M. Zani, S. Magni, G.M. Vanacore, A. Tagliaferri, J. Nucl. Mater. 393 (2009) 522–526. [23] Q. Yan, X. Zhang, J. Nucl. Mater. 444 (2014) 428–434, http://dx.doi.org/ 10.1016/j.jnucmat.2013.10.013. [24] J. Reiser, M. Rieth, B. Dafferner, A. Hoffmann, J. Nucl. Mater. 442 (2013) S204– S207. [25] R. Liu, Y. Zhou, T. Hao, T. Zhang, X.P. Wang, C.S. Liu, Q.F. Fang, J. Nucl. Mater. 424 (2012) 171–175. [26] M. Rieth, A. Hoffmann, Int. J. Refract. Met. Hard Mater. 28 (2010) 679–686. [27] T. Hirai, G. Pintsuk, J. Linke, M. Batilliot, J. Nucl. Mater. 390 (2009) 751–754.
Contents lists available at ScienceDirect
Journal of Nuclear Materials journal homepage: www.elsevier.com/locate/jnucmat
Morphology evolution of La2O3 and crack characteristic in W–La2O3 alloy under transient heat loading Xiaoxin Zhang, Qingzhi Yan ⇑ Institute of Nuclear Materials, University of Science & Technology Beijing, 30 Xueyuan Road, Haidian District, Beijing, China
h i g h l i g h t s 2
Thermal shock tests were performed on un-heated WL10 samples at 0.22–1.1 GW/m . 2
The surfaces loaded at 0.88 and 1.1 GW/m were divided into four areas. In the transition area, the left vacancies aggravated the crack initiation. In the resolidification area, the craters did not facilitate the crack initiation. WL10 showed worse performance than pure tungsten when exposed to thermal loads.
a r t i c l e
i n f o
Article history: Received 21 October 2013 Accepted 1 April 2014 Available online 13 April 2014
a b s t r a c t Thermal shock tests were performed on W–La2O3 specimens by an electron beam to investigate the morphology evolution of La2O3 and crack characteristics. Cracks appeared at power densities ranging from 0.22 to 0.88 GW/m2 and disappeared completely at 1.1 GW/m2. The specimens tested at 0.22–0.66 GW/m2 kept their original morphologies except for cracks. For the specimens loaded at 0.88 and 1.1 GW/m2, the sample surfaces were divided into four areas including no heat-loaded area, two transition areas and resolidification area. In the transition areas, the left vacancies after melting, evaporation and dissociation of La2O3 particles kept their original elongated shape and aggravated the formation of crack. In the resolidification area, the craters were partly or completely filled by tungsten and the round craters with smooth edges did not facilitate the crack initiation. Thermal fatigue tests suggested that WL10 showed worse performance than pure tungsten when exposed to thermal loads. Ó 2014 Elsevier B.V. All rights reserved.
1. Introduction Tungsten (W) is considered as one of the most promising candidate for plasma facing materials (PFMs) in future fusion reactors because of its high melting point, high thermal conductivity, low deuterium/tritium retention and low sputter rates [1–4]. PFM will be subjected to heavy thermal loads: quasi-stationary heat fluxes, edge localized modes (ELMs), vertical displacement events and plasma disruptions [5–8]. W–La2O3 have been proved to exhibit better mechanical properties, higher recrystallization temperature, more toughness than pure tungsten [9–14]. So W–La2O3 is generally considered as structural material in the plasma facing component, while pure tungsten is considered as a plasma facing material, which will suffer high thermal loads [15]. To qualify the resistance to high thermal loads of tungsten, high heat load tests have been performed in plasma simulators, neutral beam facility ⇑ Corresponding author. Tel./fax: +86 01062334951. E-mail address: [email protected] (Q. Yan). http://dx.doi.org/10.1016/j.jnucmat.2014.04.001 0022-3115/Ó 2014 Elsevier B.V. All rights reserved.
and electron beam facility [16–22]. In Ref. [16], mock-ups armored with W rods (pure W and lanthanated W) embedded in watercooled copper-alloy heat sinks were prepared, and these mockups exhibited excellent performance in thermal response tests at up to 30 MW/m2 and in thermal cycling tests of 500 cycles (10 s on, 10 s off) at 25 MW/m2. In Ref. [17], the integral bonding of tungsten tube with an internal 38-mm-long tungsten porous mesh was demonstrated and this integrally bonded tungsten heat sink displayed outstanding heat removal with a peak absorbed heat flux of 22.4 MW/m2 with He at 4 MPa, flowing at 27 g/s. Especially, Refs. [18–22] showed the different performance of W–La2O3 under high heat flux tests. In Ref. [18], W–La2O3 samples were irradiated by powerful deuterium plasma shots and the doped La2O3 particles decrease the development of a mesh of microcracks on the irradiated surface, which was conditioned by the greater brittleness of pure tungsten in comparison with W–La2O3; In Ref. [19], heat loads were applied to pure tungsten and W–La2O3 samples in the neutral beam facility GLADIS. Bubble formation was greatly suppressed in W–La2O3 than pure tungsten,
284
X. Zhang, Q. Yan / Journal of Nuclear Materials 451 (2014) 283–291
Fig. 1. Microstructures of the rolled WL10 on cross-section surface (a–c) and the illustration of electron beam orientation with respect to the loaded surface and rolling direction (d), RD, TD and ND stand for rolling direction, transverse direction and normal direction, respectively.
Absorbed energy, J
10
8
6
4
2
600
700
Fig. 2. Surface temperatures of WL10 during thermal shock and fatigue tests.
800
900
1000
1100
1200
1300
Temperature, K Fig. 3. Results of the Charpy impact tests for WL10.
which was attributed to the reduced boiling of the molten tungsten by the lower melting threshold of W–La2O3 compared with pure tungsten and the consumption of the required latent heat of the evaporated La2O3. However, some opposite results reported for the thermal shock resistance of W–La2O3. In Ref. [20], behavior of W–La2O3 under plasma disruption simulation conditions was investigated in 2MK-200 facility and the wave surface structure was caused by the low melting temperature of the La2O3 particles; In Ref. [21], thermal shock tests were performed on W–La2O3 specimens with absorbed power density up to 1.1 GW/m2 for pulse
duration of 4.6 ms and significant release of LaO molecules, melting of the surfaces were clearly detected at 0.66 GW/m2 and 1.1 GW/m2, respectively. Besides, severe cracking was observed at the vicinity of the melting spots. So the authors indicated that adding La2O3 in tungsten did not mean the improvement of the thermal shock resistance and it caused introduction of additional impurity in vacuum chamber. However, the morphological evolution of La2O3 particles was not described in this study. Ghezzi
Table 1 Thermal–mechanical properties and detailed chemical impurities information of the annealed WL10 sample. Sample
Relative density (%)
WL10
97.2 452 1312 140.1 Impurities (ppm): C < 20, S < 20, Fe < 15, Ni < 10, Al < 15, Cr < 10, Cu < 10, Ca < 20, Mg < 10, Bi < 1, Sn < 5, Si < 10, K < 10, Pb < 1, As < 10
Micro-hardness (HV)
Bending strength (MPa)
Thermal conductivity (W/mK)
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Fig. 4. Photos of WL10 after the Charpy impact tests.
Fig. 5. An overview of the thermal shock response of the rolled WL10 exposed to the electron beam.
et al. examined the surface and bulk modification of W–La2O3 exposed to 100 ELM’s transient loads at the QSPA facility, but they more focused on the surface modification of tungsten as well as the evolution of La2O3 particle density from the bulk to the surface [22]. In summary, the morphological evolution of La2O3 particles under transient heat loading has not been reported and the effect of doping La2O3 on the thermal load resistance of tungsten is not completely clear by now. In present work, we focused on the morphological evolution of La2O3 particles as well as other thermal loading responses of tungsten in terms of crack characteristics and surface roughness.
2. Experimental The investigated W-1 wt% La2O3 (WL10) was prepared by powder metallurgical route followed by rolling process. The sintered billets were rolled from 25-mm-thick to 12 mm as the rolling temperature changed from the initial 1873 K to the target 1573 K with the engineering strain of 52%. The rolled samples were annealed at 1373 K in hydrogen atmosphere for 2 h to relax residual stress. Then the samples were cut into the size of 12 12 3 mm3 and polished to remove cutting induced scratches and other defects. Finally, the prepared specimens were mounted into a copper sample holder, with a heat removal system on the
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Fig. 6. SEM images of WL10 samples tested at 0.88 and 1.1 GW/m2.
Fig. 7. High magnification SEM images of slightly melted La2O3 in area B.
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Fig. 8. High magnification SEM images of the heavily melted La2O3 in area C.
reverse. The grain orientation was parallel to the heat transfer direction (see Fig. 1(d)), which corresponds to the ITER specification. Thermal shock tests were performed at absorbed power densities varied from 0.22 to 1.1 GW/m2 at an increment of 0.22 GW/ m2 with just 1 pulse. The detailed parameters for the transient high heat load tests were described in Ref. [23]. A homogeneous heat load distribution in the 4 4 mm2 beam spot was achieved by fast scanning (10 kHz) of the electron beam. Fig. 2 shows the surface temperature evolution of WL10 during thermal shock and fatigue tests. The surface temperature showed a general increase with the absorbed power density except 0.88 GW/m2 during the thermal shock tests, which should be caused by the melting of La2O3 particles. For the thermal fatigue tests, surface temperatures were 600–800 K. For Charpy impact tests, specimens with dimensions of 27 4 3 mm (length width height), a notch depth of 1 mm and a notch root radius of 0.1 mm were prepared. The notch orientation was R–N (‘‘R’’ stands for the rolling direction; ‘‘N’’ stands for the direction of the thickness of the plate. The first letter indicates the direction perpendicular to the expected crack plane while the
second letter stands for the expected direction of crack growth). Reiser also reported that this orientation resulted in optimum Charpy properties [24]. The specimens were heated to 673– 1273 K in argon atmosphere, then pushed to the support outside the furnace and hit by a striker (25 J) within about 2 s. The temperature drop was around 50 K measured by an infrared thermometer (HT-8865, China). Microstructure obversion and surface analysis were carried out with an optical microscope and scanning electron microscopy (SEM). Vickers micro-hardness tests were performed on the polished surface under a load of 1.96 N for 10 s by a MH-6 micro-hardness instrument. Three point bending (3PB) specimens of RD(37) TD(4) ND(3) were produced to measure the fracture strength. The values in the brackets stand for the dimensions along different directions. The 3PB tests were performed at room temperature with the deformation speed of 0.5 mm/min. Surface roughness was measured along five paths over the loaded surface by a laser profilometry. The thermal conductivity was calculated by the following formula:
k ¼ q cp a
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Fig. 9. SEM images of left vacancies in the resolidification area, low magnification image (a) and high magnification images (b–e) corresponding to the selected square region in Fig. 9 (a).
where k is the thermal conductivity, q is bulk density, cp is specific heat, a is the thermal diffusivity. The specific heat was measured by a thermal analyzing apparatus (Dupont 1090B, America). The thermal diffusivity was determined using a laser flash apparatus (LFA 457, Germany). Density was measured using Archimedes’ method.
3. Results and discussion Microstructures of the annealed WL10 are shown in Fig. 1. The pancake-like grains are flattened parallel to the plate surfaces (Fig. 1(a)), the needle-shaped grains are extremely elongated along the rolling direction with an aspect ratio of roughly 3.5:1 (Fig. 1(c)). Table 1 presents the thermal–mechanical properties and detailed chemical impurities information of the annealed WL10 sample. The thermal conductivity of 140.1 W/mK is higher than 121 W/
mK reported in Ref. [25], which should be attributed to the higher relative density of 97.2% compared with 95.0% in Ref. [25]. Fig. 3 shows the results of the Charpy impact tests and the photos of WL10 samples after the Charpy impact tests are presented in Fig. 4. WL10 specimens showed delamination (fractures propagate parallel to the specimen’s long side and perpendicular to the notch) at temperatures ranging from 773 K to 1273 K. The delamination of WL10 in Charpy test was also observed in Ref. [26]. According to the Charpy test results, the brittle-to-delamination transition temperature (defined in analogy to the ductile-to-brittle transition temperature (DBTT)) was 723–773 K for WL10. Thermal shock tests were carried out on un-heated WL10 samples at absorbed power densities varied from 0.22 to 1.1 GW/m2 in a step of 0.22 GW/m2 with grains elongated orientation paralleling to the heat flux. An overview of the thermal shock response of the annealed WL10 exposed to electron beam is given in Fig. 5. Cracks
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occurred in the samples loaded at power densities of 0.22– 0.88 GW/m2 and disappeared completely in the sample tested at the highest power density of 1.1 GW/m2. Furthermore, cracks displayed circular profile and branched at lower power density including 0.22, 0.44, 0.66 GW/m2. The reasons for the circular crack profile were described in Ref. [23]. But for the sample loaded at 0.88 GW/m2, the phenomenon of circular cracks branching was not detected. In other words, the crack density decreased as the power density increased. A similar phenomenon has been reported in Ref. [27], in which the microcrack networks were replaced by discontinuing networks and also disappear completely at higher power densities. Firstly, we investigated the amazing samples loaded at 0.88 GW/m2 and 1.1 GW/m2. Melting of tungsten matrix in samples loaded above 0.88 GW/m2 was detected obviously. Fig. 6 shows the SEM images of WL10 samples tested at 0.88 and 1.1 GW/m2. Basing on the distance from the center of scanning area, there are four different areas on the surface: no heat-loaded area (marked as A), two transition areas (marked as B and C), resolidification area (marked as D) with a diameter of 1–1.5 mm. In area A, the surface was not scanned by electron beam and kept the original morphology before thermal shock tests; In area B, C, the surface has been scanned homogeneously by electron beam but B area exhibited more brightness compared with C area; D represents the center of electron scanning area, the molten tungsten matrix solidified obviously. Fig. 6(c), (d), (e), (f) are high magnification images corresponding to A, B, C, D areas in Fig. 6(a) respectively. A area almostly kept the morphology before tests except cracks (Fig. 6(c)), which was caused by the temperature gradients and the related thermal stresses. In B area, La2O3 particles slightly melted and kept its original elongated shape (Figs. 6(d) and 7). Fig. 7 shows the high magnification SEM images of slightly melted La2O3 in area B, and tungsten element, oxygen element and lanthanum element were not detected resulting from the depth of left vacancies after the evaporation, dissociation of La2O3 particles went beyond the measurement range of SEM. The dissociation of La2O3 particles at 2773 K during thermal shock process has been proved in Ref. [21] and the evaporization of La2O3 particles must occured in samples loaded at 0.66, 0.88 and 1.1 GW/m2 with the surface temperatures over the melting point of La2O3 (2573 K). In area C, La2O3 particles melted heavily, then the melted La2O3 particles overflowed and covered the near tungsten matrix (Figs. 6(e) and 8). Fig. 8 shows the high magnification SEM images of the melted La2O3 in transition area corresponding to the selected square region in Fig. 6(e). The overflowed La2O3 particles were darker than tungsten matrix, which results in more brightness of area B compared with area C. In the resolidification area D, tungsten matrix melted obviously, then solidified and formed hill-like
structure (Figs. 6(f) and 9). Fig. 9 shows the high magnification SEM images of the resolidification area. Besides, the elongated left vacancies were replaced by the craters, which should be caused by the following process: (1) La2O3 particles as well as the left vacancies were covered with the molten tungsten, (2) gases originated from the evaporation, dissociation of La2O3 particles escaped into the ambient resulting in the craters formation, and (3) the craters were partly or completely filled with the molten tungsten (Fig. 9). More importantly, the left vacancies after melting, evaporation and dissociation of La2O3 particles influenced the crack formation significantly. The elongated left vacancies in the transition areas with tip aggravated crack formation as the stress concentration during the cooling phase (see Fig. 10). In the resolidification area, the craters with smooth edges did not facilitate the crack initiation. Secondly, the surface roughness and the crack characteristics in terms of crack depth and crack width were investigated. Fig. 11 shows the arithmetic mean roughness (Ra) as a function of power density. The surface roughness of loaded area increases slightly to a certain value from 0.22 to 0.44 GW/m2, stays constant (about 140 nm) from 0.44 to 0.66 GW/m2 and finally increase to a maximum value (about 320 nm) at 1.1 GW/m2. In the resolidification area, the hills developed obviously and the corresponding roughness values were about 6500 and 1300 nm for the samples loaded at 0.88 and 1.1 GW/m2 respectively. Besides, the maximal crack width and mean crack depth were quantified from the microstructures and depicted in Fig. 12. The maximal crack width (around
Fig. 10. SEM image of crack initiation in transition area.
Fig. 12. Power density dependence of crack parameters.
Fig. 11. Arithmetic mean roughness as a function of power density.
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Fig. 13. SEM images of the rolled WL10 after thermal fatigue tests.
8 lm) and depth (around 900 lm) were observed at 0.66 and 0.88 GW/m2, respectively. In the range from 0.22 to 0.88 GW/m2, both crack parameters showed a general increase with power density caused by the increase of the temperature gradients and the related thermal stresses as well as the elongated vacancies after melting, evaporation and dissociation of La2O3 particles. As the power density increases above 0.66 GW/m2, both crack parameters tended to minish and the cracks even disappeared at the sample loaded at 1.1 GW/m2, which should be caused by the decrease of the elastic modulus at the elevated temperature and the degradation of mechanical properties. Finally, thermal fatigue tests were performed on WL10 samples at an absorbed power density of 0.17 GW/m2 with pulse up to 1000. No deterioration was observed after 1 pulse which indicated the cracking threshold for WL10 was in the range of 0.17–0.22 GW/ m2. An overview of the thermal fatigue response of the rolled WL10 exposed to electron beam with 100, 1000 pulses is given in Fig. 13(a) and (b). After 100 pulses, the cracks with the maximal width of 3 lm appeared and the La2O3 particles kept their original shape without deterioration (Fig. 13(c)). After 1000 pulses, the maximal crack width increased up to 9 lm. Besides, the roughening of La2O3 particles was detected (Fig. 13(d)), which should be caused by the evaporation of La2O3 particles. In comparison with the responses of pure tungsten in Ref. [23], WL10 samples showed a worse performance than pure tungsten when exposed to thermal loads. The different responses of tungsten grades including pure tungsten and WL10 as well as potassium doped tungsten under transient heat loading will be explained in future. 4. Conclusions Thermal shock and fatigue tests were performed on specimens cut from a rolled WL10 in an electron beam facility and the grain orientation was parallel to the heat transfer direction. The brittle-to-delamination transition temperature was 723–773 K for WL10. Thermal shock tests were performed at absorbed power densities varied from 0.22 to 1.1 GW/m2 with an increment of 0.22 GW/m2. Cracks appeared in the samples loaded at power densities ranging from 0.22 to 0.88 GW/m2 and disappeared com-
pletely in the sample tested at the highest power density of 1.1 GW/m2. The specimens tested at 0.22–0.66 GW/m2 kept their original morphologies except for cracks. For the specimens loaded at 0.88 and 1.1 GW/m2, the sample surfaces were divided into four areas including no heat-loaded area, two transition areas and resolidification area. In the transition areas, the melted La2O3 overflowed, covered the near tungsten matrix and the elongated left vacancies aggravated the formation of crack as the stress concentration during the cooling phase. In the resolidification area, La2O3 particles and tungsten matrix suffered the greatest melt. The elongated left vacancies were replaced by the craters and the craters were partly or completely filled with the molten tungsten resulting in smooth edges, which did not facilitate the crack initiation. Finally, thermal fatigue tests were performed at absorbed power density of 0.17 GW/m2 with pulse up to 1000, the maximal crack widths were 3 lm and 9 lm when the pulses increased from 100 to 1000. All the findings suggested that WL10 showed a worse performance than pure tungsten when exposed to thermal loads. Acknowledgement The authors gratefully acknowledge the financial support of ITER-National Magnetic Confinement Fusion Program (2014GB123000). References [1] M. Rieth et al., J. Nucl. Mater. 417 (2011) 463–467. [2] V. Philipps, J. Nucl. Mater. 415 (2011) S2–S9. [3] T. Hao, Z.Q. Fan, S.X. Zhao, G.N. Luo, C.S. Liu, Q.F. Fang, J. Nucl. Mater. 433 (2012) 351–356. [4] Z. Zhou, J. Tan, D. Qu, G. Pintsuk, M. RöDig, J. Linke, J. Nucl. Mater. 431 (2011) 202–205. [5] M. Wirtz, J. Linke, G. Pintsuk, L. Singheiser, M. Zlobinski, J. Nucl. Mater. 438 (2013) S833–S836, http://dx.doi.org/10.1016/j.jnucmat.2013.01.180. [6] G. Pintsuk, A. Prokhodtseva, I. Uytdenhouwen, J. Nucl. Mater. 417 (2011) 481– 486. [7] M. Merola, M. RöDig, J. Linke, R. Duwe, G. Vieider, J. Nucl. Mater. 258 (1998) 653–657. [8] M. Roedig, I. Kupriyanov, J. Linke, X. Liu, Z. Wang, J. Nucl. Mater. 417 (2011) 761–764. [9] Q. Yan, X. Zhang, T. Wang, C. Yang, C. Ge, J. Nucl. Mater. 442 (2013) S233–S236, http://dx.doi.org/10.1016/j.jnucmat.2013.01.307.
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