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Impact of the surface quality on the thermal shock performance of beryllium armor tiles for first wall applications
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Impact of the surface quality on the thermal shock performance of beryllium armor tiles for first wall applications B. Spilker ∗ , J. Linke, G. Pintsuk, M. Wirtz Forschungszentrum Jülich, Institut für Energie- und Klimaforschung, 52425 Jülich, Germany
h i g h l i g h t s • • • • •
Different surface qualities of S-65 beryllium are tested under high heat flux conditions. After 1000 thermal shocks, the loaded area exhibits a crucial destruction. Stress accelerated grain boundary oxidation/dynamic embrittlement effects are linked to the thermal shock performance of beryllium. Thermally induced cracks form between 1 and 10 pulses and grow wider and deeper between 10 and 100 pulses. Thermally induced cracks form and propagate independently from surface grooves and the surface quality.
a r t i c l e
i n f o
Article history: Received 24 July 2015 Received in revised form 19 October 2015 Accepted 26 October 2015 Available online xxx Keywords: First wall Beryllium High heat flux testing ELMs Surface quality
a b s t r a c t Beryllium will be applied as first wall armor material in ITER. The armor has to sustain high steady state and transient power fluxes. For transient events like edge localized modes, these transient power fluxes rise up to 1.0 GW m−2 with a duration of 0.5–0.75 ms in the divertor region and a significant fraction of this power flux is deposited on the first wall as well. In the present work, the reference beryllium grade for the ITER first wall application S-65 was prepared with various surface conditions and subjected to transient power fluxes (thermal shocks) with ITER relevant loading parameters. After 1000 thermal shocks, a crucial destruction of the entire loaded area was observed and linked to the stress accelerated grain boundary oxidation (SAGBO)/dynamic embrittlement (DE) effect. Furthermore, the study revealed that the majority of the thermally induced cracks formed between 1 and 10 pulses and then grew wider and deeper with increasing pulse number. The surface quality did not influence the cracking behavior of beryllium in any detectable way. However, the polished surface demonstrated the highest resistance against the observed crucial destruction mechanism. © 2015 Elsevier B.V. All rights reserved.
1. Introduction The first wall (FW) in the next step experimental fusion reactor ITER will be armored with ∼620 m2 of beryllium. This armor has to sustain high steady state heat fluxes during the startup phase of the plasma and additionally the impact of transient events (thermal shock events) such as edge localized modes (ELMs), disruptions, massive gas injections, and vertical displacement events (VDEs). ELMs deposit power densities of up to 1 GW m−2 for a pulse duration around 0.5–0.75 ms on the plasma facing materials. Power loads in this order of magnitude have various effects on the materials surface, which range from roughening over crack formation to melt layer formation in the case of beryllium. This irreversible
∗ Corresponding author. E-mail address: [email protected] (B. Spilker).
damage leads to a decrease of the lifetime of the plasma facing components and needs to be reliably estimated under the expected operational conditions. Previous research concerning thermal shock loads on beryllium has focused on its behavior under disruption and VDE loading [1,2]. Additionally, work has been done on the erosion behavior under ELM-like loading of beryllium [3–5]. However, there are indications that the surface quality of the plasma facing materials has a significant influence on their thermal shock performance [6]. The FW beryllium armor tiles are cut via electrical discharge machining (EDM) and joined to the supporting heat sink structure via hot isostatic pressing (HIP). Additional grinding or polishing steps could be added to the fabrication routine, if they were justified by an improved performance and moderate cost. However, apart from the possibly gained improved thermal shock performance, the stressing of the joint and a possible decrease of its quality due to the additional machining process need to be taken into account as well.
http://dx.doi.org/10.1016/j.fusengdes.2015.10.028 0920-3796/© 2015 Elsevier B.V. All rights reserved.
Please cite this article in press as: B. Spilker, et al., Impact of the surface quality on the thermal shock performance of beryllium armor tiles for first wall applications, Fusion Eng. Des. (2015), http://dx.doi.org/10.1016/j.fusengdes.2015.10.028
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2 Table 1 Overview of the tested surface qualities. Surface quality
Polished
Ultra
Fine
Medium
Rough
EDM
Polishing/grinding particle diameter (m) Ra (m)
1 0.1
26 0.9
46 1.1
82 1.3
201 1.9
– 2.3
This work investigated the thermal shock performance of beryllium fulfilling the ITER specifications with various surface quality conditions. The electron beam facility JUDITH 1 [7] was used to simulate thermal shocks with ELM-like parameters. In addition, a series of thermal shocks with an equal absorbed power density and duration, but with varying pulse numbers, revealed the emergence and growth of a thermally induced crack network.
2. Experimental setup The investigated material was the ITER reference beryllium grade S-65 (revision E), provided by Materion Brush Inc. This grade is produced by powder metallurgy and subsequent uniaxial hot pressing. The beryllium oxide (BeO) content remains less than 0.6 wt%, leading to a purity of 99.4 wt% beryllium. The uniaxial production routine causes a microstructure with elongated grains (grain elongation ratio 1:1.8). The samples are cut from the rod via EDM in a particular cutting pattern to receive transversal oriented grains. Thus, the grains are elongated in the direction of thermal loading and heat propagation. The average grain circular equivalent diameter at the top surface is 13 m. After the final EDM cut (lathe cut is also feasible but induces strong residual stresses close to the surface) of the specimens, several surface treatments were applied, such as grinding with differently coarse SiC papers and polishing with a 1 m particle diamond suspension. Table 1 provides an overview of the prepared surface conditions. The surface qualities “ultra”, “fine”, “medium”, and “rough” refer to SiC ground surfaces with the respective increasing SiC particle sizes. The prepared surface conditions cover a broad range. This is owed to the fact that the different domestic agencies which procure the FW panels might apply different machining and joining techniques, thus, leading to varying final surface conditions. The arithmetic mean roughness Ra (average absolute deviation of the average profile) was measured before and after the thermal shock loading via laser profilometry. Thermal shocks with ITER ELM-like parameters were applied by the electron beam facility JUDITH 1. More precisely, the absorbed power density (Labs ) for all loaded areas was set to 0.9 GW m−2 , considering an electron absorption coefficient of 0.99 for beryllium, which was determined by a Monte-Carlo simulation with 120 keV electrons penetrating beryllium. The pulse duration (t) was set to 1 ms, since this is the lower machine boundary of JUDITH 1 for the pulse duration. The electron beam was scanning the loaded area with frequencies of 40 kHz and 31 kHz in the x- and y-direction, respectively. The acceleration voltage was set to 120 kV. Each loaded area was 4 mm × 4 mm and the samples were kept at room temperature (RT) during the exposure. A higher base temperature (∼250 ◦ C are foreseen as the operational temperature for the ITER FW) would have required a lower Labs than 0.9 GW m−2 to stay below the beryllium melting threshold. However, the rather high Labs of 0.9 GW m−2 was chosen to induce a clearly pronounced surface damage without melting and, therefore, to enable the comparison in thermal shock performance of the different surface qualities. Furthermore, it has been found that the base temperature of beryllium does not influence its threshold for thermal shock induced crack formation up to 250 ◦ C [1]. The loading frequency of 0.5 Hz was sufficiently low to allow a complete cool-down of the samples to the base temperature between
two pulses. The vacuum conditions in JUDITH 1 yielded an oxygen partial pressure of about 2 × 10−5 mbar. Subsequently, the specimens were analyzed by means of scanning electron microscopy (SEM), energy dispersive X-ray spectroscopy (EDX), and metallography. 3. Results and discussion A variety of surface morphology changes and thermally induced damages was observed on all loaded areas. These damages were manifested in the form of roughening, the formation of cracks and crack networks, and ultimately the formation of molten layers/parts. Furthermore, after 1000 pulses all loaded areas exhibit a crucial destruction that cannot be explained by thermally induced plastic deformation only. Fig. 1 shows a comparison of the loaded areas after 1000 pulses for three different surface qualities. The polished surface [Fig. 1(a)] demonstrated the highest resistance against this crucial destruction mechanism. Only at the edges and the center of the loaded area, which were more heavily loaded due to the electron beam scanning pattern, the material started to form beads that lift up from the surface. The EDM cut surface [Fig. 1(b)] exhibited a more strongly pronounced destruction of the surface but there was still a large fraction of the loaded area left, where only roughening and a crack network were present. Finally, the surface quality categorized as “ultra” [Fig. 1(c)] featured a crucial destruction of the entire loaded area. This extent of destruction was representative for all surface qualities that were prepared via grinding, i.e. “ultra”, “fine”, “medium”, and “rough”. Possible explanations for the crucial destruction mechanism are the stress accelerated grain boundary oxidation (SAGBO) effect [8] and the dynamic embrittlement (DE) effect [9]. These effects were observed for Ni-base superalloys and Cu-Be alloys that exhibited an accelerated cracking behavior under tensile stresses in air and oxygen rich atmospheres in contrast to vacuum or inert gas atmospheres [8,9]. The grain boundaries were embrittled by oxygen and the propagation of intergranular cracks was accelerated. Despite the fact that the present experiments were performed under vacuum conditions, the high affinity of beryllium for oxygen could have enabled the SAGBO/DE effect by using the residual oxygen from the JUDITH 1 vacuum chamber. A similar destruction behavior after 1000 pulses with 1.5 MJ m−2 absorbed energy density and t = 5 ms on beryllium at 250 ◦ C base temperature was observed but not linked to the oxidation of beryllium in [10]. The thermal shocks in the present experiment yielded a peak surface temperature close to the melting point of beryllium (1287 ◦ C), at least at the edges and the center of the loaded area. Thus, as soon as cracks started to propagate parallel to the loaded surface, thermal barriers arose and the affected parts overheated further and melted during each thermal shock. We assume that the reconnection of the molten beryllium to the bulk was prevented by a layer of BeO on both, the underlying material and the lifted up material, which was supported by EDX measurements. The melting threshold of BeO (2575 ◦ C) left a wide temperature range, in which beryllium could be already molten but covered by a solid layer of BeO. In the most extreme case, the described mechanism separated a whole macroscopic bead of beryllium covered with BeO, as it can be seen in the center of the loaded area in Fig. 1(b) and on certain spots in Fig. 1(c). An EDX measurement on such a bead
Please cite this article in press as: B. Spilker, et al., Impact of the surface quality on the thermal shock performance of beryllium armor tiles for first wall applications, Fusion Eng. Des. (2015), http://dx.doi.org/10.1016/j.fusengdes.2015.10.028
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Fig. 1. SEM overview images of the loaded areas after 1000 pulses for different surface qualities. Labs = 0.9 GW m−2 , t = 1 ms, RT. (a) Polished. (b) EDM. (c) Ultra.
Fig. 2. SEM images showing the crack/damage evolution from 10 to 1000 pulses for the surface quality “medium”. Labs = 0.9 GW m−2 , t = 1 ms, RT. (a) 10 pulses. (b) 100 pulses. (c) 1000 pulses.
structure in Fig. 1(c) revealed that the thickness of the BeO layer was at least 0.5 m. The obtained EDX signal completely matched with the expected signal of bulk BeO, thus, the minimum thickness of the BeO layer was determined by the consideration of the penetration depth of the EDX electrons (5 keV) in BeO, which was 0.5 m, calculated via Monte Carlo simulations. The oxidation rate of beryllium strongly increases above ∼800 ◦ C [11] and is therefore influenced by the peak temperature during the thermal shocks, but also by the total duration of the material at a temperature above ∼800 ◦ C, which was several ms for each thermal shock. However, even for 1000 pulses the total time above the oxidation temperature was only in the range of several seconds. A possible explanation for the difference of the resistance against the crucial destruction mechanism between the polished and the ground surfaces could be found in the different BeO content at the surface that was present before the thermal shock loading. The EDX analyses of the unloaded surfaces showed that there was almost no BeO present at the polished surface (just the passivating layer of a few nm), while the BeO signal of the ground surfaces was significantly higher (by a factor of ∼2). The effective beryllium surface was strongly enlarged by the wavy structure that the grinding process induced. Thus, the amount of beryllium that could be oxidized near the surface was accordingly greater. A greater amount of BeO present at the surface could accelerate the crack propagation and the separation of molten beryllium from the bulk during the thermal shock loading. However, the better performance of the EDM cut surface compared to the ground surfaces could be possibly explained by the rather planar surface structure that the EDM cut induced. The unloaded EDM cut surface exhibited many small cracks and also some small particles which caused an increase of the Ra value. Nevertheless, even though the Ra value appeared to be higher for the EDM cut surface, the SEM images indicated that the surface structure was not wavy and it was not enlarged in the same way as the ground surfaces were. Thus, the effective surface area that was available for oxidation during the thermal shock loading for the EDM cut surface was larger compared to the polished surface but smaller compared to the ground surfaces.
The evolution of the crack formation and the crack network formation with dependence on the pulse number is shown in Fig. 2. After the first pulse, barely any detectable cracks had formed. For 10 pulses, the edges and the center of the loaded area showed the formation of a crack network and an increase in roughness was detectable. The crack parameters were evaluated from the SEM images and metallographic cross section images and plotted in Fig. 3 for 10–1000 pulses. The parameters could not be evaluated for 1 pulse because of the lack of measurable cracks. The crack distance (the distance between adjacent cracks) and the crack width could also not be evaluated for 1000 pulses, because it was not possible to clearly identify a typical crack on the fully destructed surfaces in the SEM images. After 100 pulses, the edges of the loaded area appeared to be partially molten. The arithmetic mean roughness increased from 10 to 100 pulses by 1.5 m. Furthermore, the crack width increased significantly after 100 pulses from ∼6 m to ∼14 m in the case of the polished surface [Fig. 3(b)] and the crack network was more distinct. Overall, the surface quality did not affect the cracking parameters as it can be seen in Fig. 3. The crack distance and the crack width were within the standard deviation in the same range for all surface qualities with the same applied pulse number. After 1000 pulses, the full destruction of the loaded area is shown in Fig. 2(c), which coincided with a strong increase of Ra to 42 m. The distribution of crack depths for all surface qualities and 10–1000 pulses is plotted in Fig. 3(c). For 10 pulses, almost all of the cracks that were visible in the cross sections had a depth below 100 m, except for a few that had a depth between 101 and 200 m. For 100 pulses, the largest fraction of the cracks still had a depth below 100 m, but many deeper cracks were detected, some of them even deeper than 500 m. This change of the crack depth distribution from 10 to 100 pulses indicated that during the first 10 pulses initial cracks with a small depth below 100 m formed and some of these initial cracks then grew deeper with increasing pulse number and acted as a natural castellation of the loaded area. After 1000 pulses, the crucial destruction of the loaded area let many of the shallow cracks disappear in the cross section within the depth of the destructed layer. In addition, some
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Fig. 3. Crack parameters for 10–1000 pulses and for all tested surface conditions. Labs = 0.9 GW m−2 , t = 1 ms, RT. (a) Crack distance, measured on a 200 m × 200 m lattice vertically and horizontally, averaged. (b) Crack width, averaged. (c) Crack depth histogram; the crack number within each interval is counted for each loaded area (including all surface qualities) and normalized to the total number of cracks for the respective pulse number. Total number of measurable cracks in the cross sections: 51 (10 pulses), 87 (100 pulses), 40 (1000 pulses).
Fig. 4. SEM images of representative crack initiation and propagation sites after 10 pulses for different surface qualities. Labs = 0.9 GW m−2 , t = 1 ms, RT. (a) Ultra. (b) Fine. (c) Rough.
cracks grew down to a depth of 864 m at maximum on the “rough” sample. Representative crack tips, crack initiation, and crack propagation sites are illustrated in Fig. 4 for three different surface qualities. The cracks propagated completely independent of pre-existing surface grooves of any direction and size. Even small, shallow cracks that had just formed and were expected to be more sensitive to pre-existing surface structures were not influenced by the surface grooves. This indicated that the crack evolution and cracking behavior was not influenced by the surface quality at all under the presented conditions. 4. Summary and conclusions The reference beryllium grade for the ITER FW S-65 was tested under ELM-like loading conditions with varying pulse numbers and surface qualities in the electron beam facility JUDITH 1. The investigated sample surface qualities were either prepared by polishing with a 1 m particle diamond suspension, applying SiC papers with different grinding particle sizes, or left in the as received EDM cut condition. A major observation after the exposure with 1000 thermal shocks of 0.9 GW m−2 with a pulse duration of 1 ms was the crucial destruction of the loaded area that cannot be explained by thermally induced plastic deformation only. This destruction mechanism could be driven by the SAGBO/DE effect, but further experiments with varying oxygen partial pressures and/or inert gas atmospheres are proposed to investigate the influence of residual oxygen on the thermal shock performance of beryllium. The threshold for the crucial destruction mechanism is dependent on the peak surface temperature during the thermal shocks and on the total duration (i.e. pulse duration and pulse number) of the affected material at a temperature above the beryllium oxidation threshold temperature which was found to be ∼800 ◦ C, in agreement with literature. In the present experiments, this threshold temperature was exceeded only for several ms for each thermal shock.
However, the total time of the surface near material above the oxidation temperature of beryllium was sufficient to deteriorate its thermal shock performance by the crucial destruction mechanism that was linked to the massive formation of BeO at the loaded surface. The minimum thickness of the BeO layer after 1000 pulses was found to be 0.5 m, determined by EDX measurements with the consideration of the EDX electron penetration depth, which was calculated via Monte Carlo simulations. The polished sample demonstrated the highest resistance against the crucial destruction mechanism, while all ground surfaces experienced a destruction of the entire loaded area. This behavior was linked to the pre-existing amount of BeO at the unloaded surfaces. A higher amount of BeO present at the surface accelerated the observed destruction. Furthermore, the thermally induced crack evolution was studied depending on the pulse number in the range of 1–1000 pulses. After one pulse, barely any cracks had formed. The majority of the cracks formed between 1 and 10 pulses and had a crack depth smaller than 100 m after 10 pulses. Further thermal shocks increased both, the crack width and the crack depth but the crack distance did not increase significantly, which led to the conclusion that barely any new cracks formed between 10 and 100 pulses. The crucial destruction mechanism that was observed for 1000 pulses made the analysis of thermally induced cracks difficult. Most of the shallow cracks vanished within the destructed layer but some deep cracks with a crack depth of 864 m at maximum were detected. Additional weight loss measurements could reveal, whether the material of the destructed area gets eroded during the thermal shocks or if the material changes its structure (e.g. via the observed bead-like structure formation) without net loss of material. Finally, a close investigation of representative crack initiation and propagation sites revealed that the cracks form and propagate independently of any pre-existing surface grooves or preparations. In conclusion, the crack evolution and the cracking behavior of beryllium were not influenced by the surface quality in the framework of the applied testing conditions. The present results indicate that there is no obligation for an additional surface treatment of
Please cite this article in press as: B. Spilker, et al., Impact of the surface quality on the thermal shock performance of beryllium armor tiles for first wall applications, Fusion Eng. Des. (2015), http://dx.doi.org/10.1016/j.fusengdes.2015.10.028
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the beryllium armor tiles for ITER to improve their thermal shock resistance in terms of the cracking behavior. However, the polished sample demonstrated the highest resistance against the observed crucial destruction mechanism. Nevertheless, experiments with higher pulse numbers of up to 106 are proposed to reveal whether the observed advantage of the polished surface for 1000 pulses holds true for further increasing pulse numbers. Acknowledgments The authors would like to kindly thank T. Flossdorf and Dr. E. Wessel for their benevolent assistance in preparing the samples and retrieving the SEM analyses.
References [1] [2] [3] [4] [5] [6] [7] [8]
M. Roedig, et al., J. Nucl. Mater. 417 (2011) 761–764. J. Linke, Trans. Fusion Sci. Technol. 49 (2006) 455–464. I.B. Kupriyanov, et al., Fusion Eng. Des. 89 (2014) 1074–1080. I.B. Kupriyanov, et al., J. Nucl. Mater. 463 (2015) 781–786. B. Bazylev, et al., J. Nucl. Mater. 386–388 (2009) 919–921. X. Liu, et al., J. Nucl. Mater. 442 (2013) S309–S312. R. Duwe, et al., Fusion Technol. (1995) 356–358. D.F. Smith, et al., Physical Metallurgy of Controlled Expansion Invar-Type Alloys, TMS, Warrendale, PA, 1990, pp. 253. [9] U. Krupp, Int. Mater. Rev. 50 (2005) 83–97. [10] V. Barabash, et al., Phys. Scr. T145 (014007) (2011) 6. [11] A.J. Stonehouse, et al., Mater. Protect. 4 (1) (1965) 24–28.
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ARTICLE IN PRESS
FUSION-8306; No. of Pages 5
Fusion Engineering and Design xxx (2015) xxx–xxx
Contents lists available at ScienceDirect
Fusion Engineering and Design journal homepage: www.elsevier.com/locate/fusengdes
Impact of the surface quality on the thermal shock performance of beryllium armor tiles for first wall applications B. Spilker ∗ , J. Linke, G. Pintsuk, M. Wirtz Forschungszentrum Jülich, Institut für Energie- und Klimaforschung, 52425 Jülich, Germany
h i g h l i g h t s • • • • •
Different surface qualities of S-65 beryllium are tested under high heat flux conditions. After 1000 thermal shocks, the loaded area exhibits a crucial destruction. Stress accelerated grain boundary oxidation/dynamic embrittlement effects are linked to the thermal shock performance of beryllium. Thermally induced cracks form between 1 and 10 pulses and grow wider and deeper between 10 and 100 pulses. Thermally induced cracks form and propagate independently from surface grooves and the surface quality.
a r t i c l e
i n f o
Article history: Received 24 July 2015 Received in revised form 19 October 2015 Accepted 26 October 2015 Available online xxx Keywords: First wall Beryllium High heat flux testing ELMs Surface quality
a b s t r a c t Beryllium will be applied as first wall armor material in ITER. The armor has to sustain high steady state and transient power fluxes. For transient events like edge localized modes, these transient power fluxes rise up to 1.0 GW m−2 with a duration of 0.5–0.75 ms in the divertor region and a significant fraction of this power flux is deposited on the first wall as well. In the present work, the reference beryllium grade for the ITER first wall application S-65 was prepared with various surface conditions and subjected to transient power fluxes (thermal shocks) with ITER relevant loading parameters. After 1000 thermal shocks, a crucial destruction of the entire loaded area was observed and linked to the stress accelerated grain boundary oxidation (SAGBO)/dynamic embrittlement (DE) effect. Furthermore, the study revealed that the majority of the thermally induced cracks formed between 1 and 10 pulses and then grew wider and deeper with increasing pulse number. The surface quality did not influence the cracking behavior of beryllium in any detectable way. However, the polished surface demonstrated the highest resistance against the observed crucial destruction mechanism. © 2015 Elsevier B.V. All rights reserved.
1. Introduction The first wall (FW) in the next step experimental fusion reactor ITER will be armored with ∼620 m2 of beryllium. This armor has to sustain high steady state heat fluxes during the startup phase of the plasma and additionally the impact of transient events (thermal shock events) such as edge localized modes (ELMs), disruptions, massive gas injections, and vertical displacement events (VDEs). ELMs deposit power densities of up to 1 GW m−2 for a pulse duration around 0.5–0.75 ms on the plasma facing materials. Power loads in this order of magnitude have various effects on the materials surface, which range from roughening over crack formation to melt layer formation in the case of beryllium. This irreversible
∗ Corresponding author. E-mail address: [email protected] (B. Spilker).
damage leads to a decrease of the lifetime of the plasma facing components and needs to be reliably estimated under the expected operational conditions. Previous research concerning thermal shock loads on beryllium has focused on its behavior under disruption and VDE loading [1,2]. Additionally, work has been done on the erosion behavior under ELM-like loading of beryllium [3–5]. However, there are indications that the surface quality of the plasma facing materials has a significant influence on their thermal shock performance [6]. The FW beryllium armor tiles are cut via electrical discharge machining (EDM) and joined to the supporting heat sink structure via hot isostatic pressing (HIP). Additional grinding or polishing steps could be added to the fabrication routine, if they were justified by an improved performance and moderate cost. However, apart from the possibly gained improved thermal shock performance, the stressing of the joint and a possible decrease of its quality due to the additional machining process need to be taken into account as well.
http://dx.doi.org/10.1016/j.fusengdes.2015.10.028 0920-3796/© 2015 Elsevier B.V. All rights reserved.
Please cite this article in press as: B. Spilker, et al., Impact of the surface quality on the thermal shock performance of beryllium armor tiles for first wall applications, Fusion Eng. Des. (2015), http://dx.doi.org/10.1016/j.fusengdes.2015.10.028
G Model FUSION-8306; No. of Pages 5
ARTICLE IN PRESS B. Spilker et al. / Fusion Engineering and Design xxx (2015) xxx–xxx
2 Table 1 Overview of the tested surface qualities. Surface quality
Polished
Ultra
Fine
Medium
Rough
EDM
Polishing/grinding particle diameter (m) Ra (m)
1 0.1
26 0.9
46 1.1
82 1.3
201 1.9
– 2.3
This work investigated the thermal shock performance of beryllium fulfilling the ITER specifications with various surface quality conditions. The electron beam facility JUDITH 1 [7] was used to simulate thermal shocks with ELM-like parameters. In addition, a series of thermal shocks with an equal absorbed power density and duration, but with varying pulse numbers, revealed the emergence and growth of a thermally induced crack network.
2. Experimental setup The investigated material was the ITER reference beryllium grade S-65 (revision E), provided by Materion Brush Inc. This grade is produced by powder metallurgy and subsequent uniaxial hot pressing. The beryllium oxide (BeO) content remains less than 0.6 wt%, leading to a purity of 99.4 wt% beryllium. The uniaxial production routine causes a microstructure with elongated grains (grain elongation ratio 1:1.8). The samples are cut from the rod via EDM in a particular cutting pattern to receive transversal oriented grains. Thus, the grains are elongated in the direction of thermal loading and heat propagation. The average grain circular equivalent diameter at the top surface is 13 m. After the final EDM cut (lathe cut is also feasible but induces strong residual stresses close to the surface) of the specimens, several surface treatments were applied, such as grinding with differently coarse SiC papers and polishing with a 1 m particle diamond suspension. Table 1 provides an overview of the prepared surface conditions. The surface qualities “ultra”, “fine”, “medium”, and “rough” refer to SiC ground surfaces with the respective increasing SiC particle sizes. The prepared surface conditions cover a broad range. This is owed to the fact that the different domestic agencies which procure the FW panels might apply different machining and joining techniques, thus, leading to varying final surface conditions. The arithmetic mean roughness Ra (average absolute deviation of the average profile) was measured before and after the thermal shock loading via laser profilometry. Thermal shocks with ITER ELM-like parameters were applied by the electron beam facility JUDITH 1. More precisely, the absorbed power density (Labs ) for all loaded areas was set to 0.9 GW m−2 , considering an electron absorption coefficient of 0.99 for beryllium, which was determined by a Monte-Carlo simulation with 120 keV electrons penetrating beryllium. The pulse duration (t) was set to 1 ms, since this is the lower machine boundary of JUDITH 1 for the pulse duration. The electron beam was scanning the loaded area with frequencies of 40 kHz and 31 kHz in the x- and y-direction, respectively. The acceleration voltage was set to 120 kV. Each loaded area was 4 mm × 4 mm and the samples were kept at room temperature (RT) during the exposure. A higher base temperature (∼250 ◦ C are foreseen as the operational temperature for the ITER FW) would have required a lower Labs than 0.9 GW m−2 to stay below the beryllium melting threshold. However, the rather high Labs of 0.9 GW m−2 was chosen to induce a clearly pronounced surface damage without melting and, therefore, to enable the comparison in thermal shock performance of the different surface qualities. Furthermore, it has been found that the base temperature of beryllium does not influence its threshold for thermal shock induced crack formation up to 250 ◦ C [1]. The loading frequency of 0.5 Hz was sufficiently low to allow a complete cool-down of the samples to the base temperature between
two pulses. The vacuum conditions in JUDITH 1 yielded an oxygen partial pressure of about 2 × 10−5 mbar. Subsequently, the specimens were analyzed by means of scanning electron microscopy (SEM), energy dispersive X-ray spectroscopy (EDX), and metallography. 3. Results and discussion A variety of surface morphology changes and thermally induced damages was observed on all loaded areas. These damages were manifested in the form of roughening, the formation of cracks and crack networks, and ultimately the formation of molten layers/parts. Furthermore, after 1000 pulses all loaded areas exhibit a crucial destruction that cannot be explained by thermally induced plastic deformation only. Fig. 1 shows a comparison of the loaded areas after 1000 pulses for three different surface qualities. The polished surface [Fig. 1(a)] demonstrated the highest resistance against this crucial destruction mechanism. Only at the edges and the center of the loaded area, which were more heavily loaded due to the electron beam scanning pattern, the material started to form beads that lift up from the surface. The EDM cut surface [Fig. 1(b)] exhibited a more strongly pronounced destruction of the surface but there was still a large fraction of the loaded area left, where only roughening and a crack network were present. Finally, the surface quality categorized as “ultra” [Fig. 1(c)] featured a crucial destruction of the entire loaded area. This extent of destruction was representative for all surface qualities that were prepared via grinding, i.e. “ultra”, “fine”, “medium”, and “rough”. Possible explanations for the crucial destruction mechanism are the stress accelerated grain boundary oxidation (SAGBO) effect [8] and the dynamic embrittlement (DE) effect [9]. These effects were observed for Ni-base superalloys and Cu-Be alloys that exhibited an accelerated cracking behavior under tensile stresses in air and oxygen rich atmospheres in contrast to vacuum or inert gas atmospheres [8,9]. The grain boundaries were embrittled by oxygen and the propagation of intergranular cracks was accelerated. Despite the fact that the present experiments were performed under vacuum conditions, the high affinity of beryllium for oxygen could have enabled the SAGBO/DE effect by using the residual oxygen from the JUDITH 1 vacuum chamber. A similar destruction behavior after 1000 pulses with 1.5 MJ m−2 absorbed energy density and t = 5 ms on beryllium at 250 ◦ C base temperature was observed but not linked to the oxidation of beryllium in [10]. The thermal shocks in the present experiment yielded a peak surface temperature close to the melting point of beryllium (1287 ◦ C), at least at the edges and the center of the loaded area. Thus, as soon as cracks started to propagate parallel to the loaded surface, thermal barriers arose and the affected parts overheated further and melted during each thermal shock. We assume that the reconnection of the molten beryllium to the bulk was prevented by a layer of BeO on both, the underlying material and the lifted up material, which was supported by EDX measurements. The melting threshold of BeO (2575 ◦ C) left a wide temperature range, in which beryllium could be already molten but covered by a solid layer of BeO. In the most extreme case, the described mechanism separated a whole macroscopic bead of beryllium covered with BeO, as it can be seen in the center of the loaded area in Fig. 1(b) and on certain spots in Fig. 1(c). An EDX measurement on such a bead
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Fig. 1. SEM overview images of the loaded areas after 1000 pulses for different surface qualities. Labs = 0.9 GW m−2 , t = 1 ms, RT. (a) Polished. (b) EDM. (c) Ultra.
Fig. 2. SEM images showing the crack/damage evolution from 10 to 1000 pulses for the surface quality “medium”. Labs = 0.9 GW m−2 , t = 1 ms, RT. (a) 10 pulses. (b) 100 pulses. (c) 1000 pulses.
structure in Fig. 1(c) revealed that the thickness of the BeO layer was at least 0.5 m. The obtained EDX signal completely matched with the expected signal of bulk BeO, thus, the minimum thickness of the BeO layer was determined by the consideration of the penetration depth of the EDX electrons (5 keV) in BeO, which was 0.5 m, calculated via Monte Carlo simulations. The oxidation rate of beryllium strongly increases above ∼800 ◦ C [11] and is therefore influenced by the peak temperature during the thermal shocks, but also by the total duration of the material at a temperature above ∼800 ◦ C, which was several ms for each thermal shock. However, even for 1000 pulses the total time above the oxidation temperature was only in the range of several seconds. A possible explanation for the difference of the resistance against the crucial destruction mechanism between the polished and the ground surfaces could be found in the different BeO content at the surface that was present before the thermal shock loading. The EDX analyses of the unloaded surfaces showed that there was almost no BeO present at the polished surface (just the passivating layer of a few nm), while the BeO signal of the ground surfaces was significantly higher (by a factor of ∼2). The effective beryllium surface was strongly enlarged by the wavy structure that the grinding process induced. Thus, the amount of beryllium that could be oxidized near the surface was accordingly greater. A greater amount of BeO present at the surface could accelerate the crack propagation and the separation of molten beryllium from the bulk during the thermal shock loading. However, the better performance of the EDM cut surface compared to the ground surfaces could be possibly explained by the rather planar surface structure that the EDM cut induced. The unloaded EDM cut surface exhibited many small cracks and also some small particles which caused an increase of the Ra value. Nevertheless, even though the Ra value appeared to be higher for the EDM cut surface, the SEM images indicated that the surface structure was not wavy and it was not enlarged in the same way as the ground surfaces were. Thus, the effective surface area that was available for oxidation during the thermal shock loading for the EDM cut surface was larger compared to the polished surface but smaller compared to the ground surfaces.
The evolution of the crack formation and the crack network formation with dependence on the pulse number is shown in Fig. 2. After the first pulse, barely any detectable cracks had formed. For 10 pulses, the edges and the center of the loaded area showed the formation of a crack network and an increase in roughness was detectable. The crack parameters were evaluated from the SEM images and metallographic cross section images and plotted in Fig. 3 for 10–1000 pulses. The parameters could not be evaluated for 1 pulse because of the lack of measurable cracks. The crack distance (the distance between adjacent cracks) and the crack width could also not be evaluated for 1000 pulses, because it was not possible to clearly identify a typical crack on the fully destructed surfaces in the SEM images. After 100 pulses, the edges of the loaded area appeared to be partially molten. The arithmetic mean roughness increased from 10 to 100 pulses by 1.5 m. Furthermore, the crack width increased significantly after 100 pulses from ∼6 m to ∼14 m in the case of the polished surface [Fig. 3(b)] and the crack network was more distinct. Overall, the surface quality did not affect the cracking parameters as it can be seen in Fig. 3. The crack distance and the crack width were within the standard deviation in the same range for all surface qualities with the same applied pulse number. After 1000 pulses, the full destruction of the loaded area is shown in Fig. 2(c), which coincided with a strong increase of Ra to 42 m. The distribution of crack depths for all surface qualities and 10–1000 pulses is plotted in Fig. 3(c). For 10 pulses, almost all of the cracks that were visible in the cross sections had a depth below 100 m, except for a few that had a depth between 101 and 200 m. For 100 pulses, the largest fraction of the cracks still had a depth below 100 m, but many deeper cracks were detected, some of them even deeper than 500 m. This change of the crack depth distribution from 10 to 100 pulses indicated that during the first 10 pulses initial cracks with a small depth below 100 m formed and some of these initial cracks then grew deeper with increasing pulse number and acted as a natural castellation of the loaded area. After 1000 pulses, the crucial destruction of the loaded area let many of the shallow cracks disappear in the cross section within the depth of the destructed layer. In addition, some
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Fig. 3. Crack parameters for 10–1000 pulses and for all tested surface conditions. Labs = 0.9 GW m−2 , t = 1 ms, RT. (a) Crack distance, measured on a 200 m × 200 m lattice vertically and horizontally, averaged. (b) Crack width, averaged. (c) Crack depth histogram; the crack number within each interval is counted for each loaded area (including all surface qualities) and normalized to the total number of cracks for the respective pulse number. Total number of measurable cracks in the cross sections: 51 (10 pulses), 87 (100 pulses), 40 (1000 pulses).
Fig. 4. SEM images of representative crack initiation and propagation sites after 10 pulses for different surface qualities. Labs = 0.9 GW m−2 , t = 1 ms, RT. (a) Ultra. (b) Fine. (c) Rough.
cracks grew down to a depth of 864 m at maximum on the “rough” sample. Representative crack tips, crack initiation, and crack propagation sites are illustrated in Fig. 4 for three different surface qualities. The cracks propagated completely independent of pre-existing surface grooves of any direction and size. Even small, shallow cracks that had just formed and were expected to be more sensitive to pre-existing surface structures were not influenced by the surface grooves. This indicated that the crack evolution and cracking behavior was not influenced by the surface quality at all under the presented conditions. 4. Summary and conclusions The reference beryllium grade for the ITER FW S-65 was tested under ELM-like loading conditions with varying pulse numbers and surface qualities in the electron beam facility JUDITH 1. The investigated sample surface qualities were either prepared by polishing with a 1 m particle diamond suspension, applying SiC papers with different grinding particle sizes, or left in the as received EDM cut condition. A major observation after the exposure with 1000 thermal shocks of 0.9 GW m−2 with a pulse duration of 1 ms was the crucial destruction of the loaded area that cannot be explained by thermally induced plastic deformation only. This destruction mechanism could be driven by the SAGBO/DE effect, but further experiments with varying oxygen partial pressures and/or inert gas atmospheres are proposed to investigate the influence of residual oxygen on the thermal shock performance of beryllium. The threshold for the crucial destruction mechanism is dependent on the peak surface temperature during the thermal shocks and on the total duration (i.e. pulse duration and pulse number) of the affected material at a temperature above the beryllium oxidation threshold temperature which was found to be ∼800 ◦ C, in agreement with literature. In the present experiments, this threshold temperature was exceeded only for several ms for each thermal shock.
However, the total time of the surface near material above the oxidation temperature of beryllium was sufficient to deteriorate its thermal shock performance by the crucial destruction mechanism that was linked to the massive formation of BeO at the loaded surface. The minimum thickness of the BeO layer after 1000 pulses was found to be 0.5 m, determined by EDX measurements with the consideration of the EDX electron penetration depth, which was calculated via Monte Carlo simulations. The polished sample demonstrated the highest resistance against the crucial destruction mechanism, while all ground surfaces experienced a destruction of the entire loaded area. This behavior was linked to the pre-existing amount of BeO at the unloaded surfaces. A higher amount of BeO present at the surface accelerated the observed destruction. Furthermore, the thermally induced crack evolution was studied depending on the pulse number in the range of 1–1000 pulses. After one pulse, barely any cracks had formed. The majority of the cracks formed between 1 and 10 pulses and had a crack depth smaller than 100 m after 10 pulses. Further thermal shocks increased both, the crack width and the crack depth but the crack distance did not increase significantly, which led to the conclusion that barely any new cracks formed between 10 and 100 pulses. The crucial destruction mechanism that was observed for 1000 pulses made the analysis of thermally induced cracks difficult. Most of the shallow cracks vanished within the destructed layer but some deep cracks with a crack depth of 864 m at maximum were detected. Additional weight loss measurements could reveal, whether the material of the destructed area gets eroded during the thermal shocks or if the material changes its structure (e.g. via the observed bead-like structure formation) without net loss of material. Finally, a close investigation of representative crack initiation and propagation sites revealed that the cracks form and propagate independently of any pre-existing surface grooves or preparations. In conclusion, the crack evolution and the cracking behavior of beryllium were not influenced by the surface quality in the framework of the applied testing conditions. The present results indicate that there is no obligation for an additional surface treatment of
Please cite this article in press as: B. Spilker, et al., Impact of the surface quality on the thermal shock performance of beryllium armor tiles for first wall applications, Fusion Eng. Des. (2015), http://dx.doi.org/10.1016/j.fusengdes.2015.10.028
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the beryllium armor tiles for ITER to improve their thermal shock resistance in terms of the cracking behavior. However, the polished sample demonstrated the highest resistance against the observed crucial destruction mechanism. Nevertheless, experiments with higher pulse numbers of up to 106 are proposed to reveal whether the observed advantage of the polished surface for 1000 pulses holds true for further increasing pulse numbers. Acknowledgments The authors would like to kindly thank T. Flossdorf and Dr. E. Wessel for their benevolent assistance in preparing the samples and retrieving the SEM analyses.
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