Thick boron carbide coatings for protection of tokamak first wall and divertor

Fusion Engineering and Design 45 (1999) 343 – 360 www.elsevier.com/locate/fusengdes

Review article

Thick boron carbide coatings for protection of tokamak first wall and divertor O.I. Buzhinskij *, Yu.M. Semenets Troitsk Institute for Inno6ation and Fusion Research, Troitsk, Moscow reg., 142092, Russia Received 15 October 1998; received in revised form 21 December 1998; accepted 22 January 1999

Abstract A review of characteristics of various types of boron carbide coatings considered as candidate materials for protection of tokamak inner surfaces against high energy heat fluxes is presented. Such coatings are produced by various methods: chemical vapor deposition by means of chloride and fluoride techniques, gas conversion, plasma spray and reaction-sintering. Contrary to pure carbon materials, B4C has much lower chemical and high-temperature sputtering, is capable to oxygen gettering and lower hydrogen recycling. In contrast to thin boronization films, the thick coatings can resist high heat fluxes such as in tokamak divertors. Comparative analysis shows that coatings produced by the diffusion methods, such as fluoride CVD and gas conversion, are more resistent to heat loads, and one of the most promising candidates are the fluoride CVD coatings. © 1999 Elsevier Science S.A. All rights reserved. Keywords: Boron carbide; Tokamak; First wall; Divertor; High heat flux

1. Introduction The use of materials with low Z for protection of tokamak first wall and other in vessel elements is a conventional method to decrease heavy impurity influx into tokamak plasma. Recently, graphite and carbon films produced by carbonization were the most widespread materials. How-

* Corresponding author. Tel.: +95-3-340538; fax: +95-3345776. E-mail address: [email protected] (O.I. Buzhinskij)

ever the main disadvantages of a pure carbon material are its strong chemical erosion and high temperature (T\ 1000°C) sublimation. Presently one of the most important tendencies is the use of boron-containing materials and films. Primarily the boron–carbon material uglesitall USB-15 was tested as a limiter material in the T-3 tokamak [1]. At the same time its chemical erosion was tested and appeared to be much less as compared to carbon. In situ plasma chemical vapor deposition of a thin amorphous boron–carbon film, so called boronization, was for the first time performed in the TEXTOR tokamak [2].

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As a result, chemical and high temperature sputtering were significantly reduced, carbon and oxygen impurities were suppressed, and plasma performance was improved. However, bulk boron-containing material has low thermal conductivity, and thin ( B100 nm) boronization film quickly erodes on the areas subjected to high heat flux. It seems more reasonable to use thick (100– 500 mm) boron – carbon coatings on substrates with high thermal conductivity. These coatings have higher heat load durability as compared to thin boronized films. Their B/C ratio is higher, and their thermomechanical properties (under appropriate choice of substrate) are better as compared to bulk boron-containing materials.

2. Methods of fabrication There are a number of coating fabrication methods: the chemical vapor deposition (CVD) with the use of BCl3 (chloride technique) and BF3 (fluoride technique), the gas conversion, the plasma spray and the reaction-sintering or the conversion via liquid phase.

2.1. Chloride CVD technique Coatings are deposited on a substrate surface from a gas mixture of BCl3 with methane CH4 and hydrogen H2 at high temperatures (1500°C) [3 – 7]: 4BCl3 +CH4 +4H2 =B4C+ 12HCl. C3H6 can be substituted for CH4 [5 –8]. The microstructure of the coating is a mixture of boron–carbide and graphite. The B –C ratio varies in a wide range, e.g. 1.8 in ref. [3,4], and 2.6 and 3.0 for two coatings investigated in ref. [9]. Many pores near the interface of dendritic B4C and graphite have been observed [3,4]. The main phase of the coatings obtained with methane was carbon in ref. [6,7] so that the gas mixture only with hydrogen was used: 2BCl3 +3H2 =2B +6HCl

Fig. 1. Temperature dependence of the deposition rate from boron – chloride mixture.

1500°C and drops under further temperature rise because of the reverse process of coating etching. The thermodynamical calculations of B–Cl–H system [7] show that a layer of a solid boron with a crystalline structure dependent on a substrate temperature is deposited. At temperatures 1200– 1900°C the crystalline boron weakly interacts with dense graphite of the substrate, and carbidization does not occur. The morphology of the coatings obtained at 1800°C is weakly crystalline (Fig. 2) [6,7]. The degree of crystallinity increases with increasing temperature, but adherence makes worse. The optimum regime in the chloride system without methane takes place at about 1550°C [6,7].

(1)

Fig. 1 shows the dependence of the deposition rate on temperature [6,7]. It is seen that the deposition rate reaches its maximum value at

Fig. 2. Microphoto of the coating obtained with BCl3 at T= 1800°C.

O.I. Buzhinskij, Y.M. Semenets / Fusion Engineering and Design 45 (1999) 343–360

Fig. 3. The carrying-away rate of boron-containing mixture under fluoride technique. 1, tablets of amorphous boron; 2, boron carbide; 3, boron nitride wastes.

The substrates for this technique have been isotropic graphites, 2-directional C – C composite CC 1001 G, felt-type C – C composite CC312, a 2-dimensional felt-type CFC composite, and metals.

2.2. Fluoride CVD technique A reaction similar to Eq. (1) is not valid in the temperature range from 1200 to 1900°C because of high thermostability of BF3. Therefore the method of the chemical transport reactions is used in this case. A mixture consisting of grains of crystalline boron, pellets of amorphous boron or boron nitride wastes is placed into the most hot region of a two chamber reactor where the temperature ranges from 1800 to 1900°C. Gaseous boron fluoride filters through the mixture, producing the lower fluorides by the reactions: 2BF3 + B4C =6BF +C, 8BF3 + B4C =12BF2 +C; BF3 +2B =3BF, 2BF3 + B = 3BF2; BF3 +2BN =3BF +N2, 2BF3 + BN = 3BF2 + 12N2. The carrying-away rate of a boron-containing mixture is presented in Fig. 3 [7,10]. The maximum value is reached at temperatures somewhat

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above 1700°C. The coating deposition process consists of two stages. The first one involves reactions producing high activity boron on the substrate surface: 3BF2 = B+ 2BF3, 3BF= 2B+ BF3. The second is an interaction of this boron with carbon of the substrate to form a boron carbide coating: 4B+ C= B4C. The thermodynamical calculations of the B–F– H system shows that the coatings produced by the fluoride technique have no deposit of free boron [5–7]. In this system the partial pressure of boron vapor in a gas phase is lower than the equilibrium boron vapor partial pressure at a solid boron surface but higher than that at a solid boron carbide surface, so that the crystalline boron carbide is the only solid phase possible at temperatures sufficient for the reaction of boron vapor with carbon. X-ray analysis shows that the fluoride technique coatings contain B4C–B13C2 phase of rombohedral boron carbide and carbon, whatever boron containing mixture is used [5– 7,11–13]. Morphology of the coating is greatly influenced by a substrate temperature which determines the difference in temperatures between the boron containing mixture and the deposition zone, and as a result the boron oversaturation in the gas phase [6,7,10]. If this difference is more than 300°C, oversaturation is too high and the deposited layer have porous dendritic morphology (Fig. 4). Under lower oversaturation (B 150°C) the dense finegrain crystalline coating is deposited (Fig. 5). The degree of crystallinity increases with increasing temperature (Fig. 6), and there appear large crystallizes forming pentohedral structures. At the same time adherence somewhat degrades. The optimum temperature for coating deposition on graphite materials is about 1850°C for the fluoride technique. The deposition rate is from 100 to 150 mm/h, coating thickness is in the range of 100–500 mm, and the structure is a mixture of B4C and B13C2 phases. It should be noted that, due to diffusive character of coating formation and its nonstoichiometric phases, adhesion to substrates is good and preserved under high heat fluxes whereas chloride coatings may exfoliate.

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Fig. 4. Microphoto of fluoride CVD coating cross section at temperature difference between mixture and deposition zone 300°C.

The finegrain graphite MPG-8, pyrolytic graphite, the titanium doped RGT graphite and the TS-1792 graphite were used as substrates under this technique.

2.3. Gas con6ersion technique The coating is produced by a reaction of gaseous B2O3 with carbon of the substrate 2B2O3 +7C =B4C +6CO at high temperatures (2000°C). The positive feature of this method is that boron carbide layer forms directly inside a

Fig. 6. Microstructure of fluoride CVD coating after deposition at high temperature.

substrate by means of penetration of the gas into its open pores, and exfoliation of such a layer is not occurred under high heat flux irradiation. The structure is two phase: B4C and graphite and may contain many micropores [4].The B–C ratio drops with the distance from the surface, e.g. almost linearly from 1.6 to 0 in depth of 200 mm in ref. [3], and vary in wide range up to almost stoichiometric value for pure boron carbide [14,15]. The thickness of the layer can amount to 1700 mm [4] and can be nonuniform, e.g. it ranged from 50 to 200 mm in ref. [9]. Isotropic graphites (including PD330S), a 2-dimensional felt-type CFC composite and the PCC2S felt-type C–C composite have been the substrates for this method.

2.4. Plasma spray

Fig. 5. Microphoto of fluoride CVD coating cross section at temperature difference between mixture and deposition zone 150°C.

Under this method, B4C powder (with particles of some micrometers to some tens micrometers in size) is placed into a plasma gun with the Laval nozzle. The partly melted, partly overheated (dependent on their sizes) particles impinge a substrate, flatten and form a coating. There are the

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low pressure (50 – 300 mbar) plasma spray (LPPS) [8,9,18,19] and the high pressure (up to 2 bar) plasma spray (HPPS) [20 – 23]. The plasma energy density and the deposition rate are higher under HPPS. The thickness of the coating sprayed directly on the substrate is limited because the strong difference in thermal expansion coefficients between the coating and substrate induces significant stresses at the surface and causes detachment. To avoid this phenomenon a mixture of substrate material and B4C is used to produce an intermediate layer [21,22]. The fraction of a substrate material is gradually decreases with increasing the deposition thickness thereby creating a gradual transition in thermal expansion coefficient. The X-ray diffraction analysis shows that irrespective of the substrate in all the coatings a single phase (B4C) is presented. The porosity depends on the chamber pressure and can be lowered down to 8% [8]. This method enables the renewal of the coatings in situ. Graphites (EK98, AXF 5Q, K-KARB), stainless steels (1.4311, 316L, Steel 37), molybdenum, C–C composites (CC1501GR, 2-directional CC1001G, felt-type PCC-2S and CX2002U, unidirectional MFC-1), and copper have been used as substrates for this type of coatings.

2.5. Reaction-sintering Coatings are produced by brush-on technique with the use of powdered boride and carbide refractory compounds (e.g. B4C) and an organic binder [19]. The reaction is conducted under high temperatures (2400°C) and is a combination both of substrate carbon (and a binder carbon) reacting with refractory compounds and of sintering of the boride and/or carbide particles. The coatings consist of two layers: a surface layer and a graded penetration zone in the outer portion of a substrate. The surface layer microstructure is multiphase and range from reaction-sintered structures of sintered B4C particles in an eutectic formed matrix to that of hypereutectic carbon particles in a B4C – C eutectic matrix. Because of high surface energy, the coating generally develops a nonuniform thickness and topography. 40– 70% of the coated surface has a coating of 25–50

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mm thickness, and the remaining part consists of discrete nodules or thick coating regions, some of which are several millimeters in thickness. This technique has been applied to the AXF-5Q1, ATJ and EK98 isotropic graphites.

3. Erosion The temperature dependence of the sputtering yield of a fluoride CVD coating by H + ions with an energy of 10 keV in the temperature range of 150–1500°C has been studied [5,10]. The beam current density was 1 mA/cm2 and the fluence was 1020 cm − 2. The chemical sputtering was considerably suppressed (Fig. 7): the maximum sputtering yield was 5 times lower than that of pure MPG-8 graphite used as substrate, e.g. it was 5× 10 − 3 atoms/ion at 1200°C. Any effect of radiation enhanced sublimation on the sputtering yield was not observed at temperatures below 1500°C. The erosion characteristics of gas conversion coatings on a CFC composite under 3 keV D3+ (equivalent to 1 keV D + ) irradiation in the temperature range of 180–1130°C have been investigated [3,17]. The B4C conversion was effective in suppressing both chemical sputtering and radiation enhanced sublimation of graphite. The erosion yield was lower than that of the substrate (PCC-2S) by 50–60% at T =530°C, by 30–40% at 1000–1130°C and was about 3×10 − 2 atoms/ ion. The radiation enhanced sublimation was low compared to graphite at temperatures above 900°C and depended on C–B4C ratio. Such reduction in erosion was explained by the lower carbon contents and by the fact that at temperatures when the CD4 diffusion could take place, the density of the trapped D atoms decreased significantly and therefore chemical sputtering due to deutero-carbon formation is suppressed. In fusion reactors, ion fluxes onto divertor plates are considered to be more than 1022 1/m2 s, i.e. about two orders of magnitude higher than in conventional ion sources [16]. The erosion due to chemical sputtering for this high flux cannot be predicted by extrapolating data obtained from the low flux beam experiments using conventional ion sources. The sputtering yield of the gas converted

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PD-330S graphite irradiated by a high flux neutral deuterium beam with an acceleration energy of 5 keV and fluxes from 2 ×1021 1/m2 s to 1022 1/m2 s has been investigated [16]. Erosion appeared to be lower at high fluxes than at low fluxes. The sputtering yield of the gas conversion B4C was 0.06 atoms/ion (at 700°C), i.e. 50 – 60% lower than of the PD-330S graphite and was independent of flux.

4. Hydrogen retention and recycling Hydrogen sorption, desorption, retention and recycling are of great importance in present-day tokamaks. Therefore much attention has been payed to investigations of these issues. Retention and thermal desorption of hydrogen implanted by means of a 4.5 keV beam of H3+ ions with a flux of 1015 H ions/cm2 s and a fluence of 5×1018 H ions/cm2 into the gas converted PD-330S graphite have been studied [25]. The total amount of hydrogen trapped in B4C was 1.5 times larger than in graphite. The activation energies of H2 and CH4 desorptions were considerably lower than those for graphite. The desorption spectra of both H2 and CH4 had two peaks due to detrapping of B – H and C – H bonds while the

graphite had only one peak. The activation energy of the B–H bond was lower than that of the C–H bond which in its turn was lower than the activation energy of C–H bond in the graphite. Therefore the hydrogen content of the B4C coating decreased more rapidly with temperature than in graphite. This fact was also reported in ref. [4] where the thermal desorption temperature of B4C was 250°C lower than that of graphite. Retention of deuterium in chloride CVD and gas conversion coated graphite tiles exposed to discharge shots in JT-60U has been investigated [26]. It was found to range between 6×1020 1/cm2 and 6 ×1021 1/cm2. Highest deuterium retention was observed on the outboard lower part of the first wall tiles. The discharge density rise rate decreased, which means a lower particle recycling. These improvements were in good agreement with the results obtained by boronization [14]. The absolute concentration of hydrogen isotopes in fluoride CVD coated graphite tiles of DIII-D has been studied by means of a new neutron induced elastic recoil detection (NERD) method [27,28]. The tiles were produced from the TS-1792 graphite [29] and coated by means of the fluoride CVD technique with B4C layers of 80– 180 mm in thickness. The tiles were placed along

Fig. 7. Temperature dependence of the sputtering yield of the B4C coating on MPG graphite and uglesitall USB-15 under 10 keV hydrogen ions irradiation.

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the major radius direction. After 16 weeks’ exposure in 1993, they were annealed at 1000°C. It was found that redeposited carbon layers had been formed on the tile surfaces. Their thicknesses ranged from 500 A, on the outboard tile to 5000 A, on the inboard tile. The measurements showed that hydrogen sorption took place in the upper layers of the redeposited carbon rather than in the B4C layer. Hydrogen concentration was about (1.4− 2.8)× 1018 cm − 2. The maximum sorption was observed in the tile with the maximum thickness of a redeposited carbon. Similar hydrogen concentration was observed in an uncoated tile. Deuterium and tritium were not found. Hydrogen isotope retention of fluoride CVD coated high thermal diffusivity RGT graphite [30] tile exposed to the divertor strike point of DIII-D has been studied [31,32]. The hydrogen retention inside and outside the damaged spot of the sample was 9.3× 1018 cm − 2 and 3×1017 cm − 2 respectively, and the deuterium concentration was 1.5× 1018 cm − 2 and 2× 1017 cm − 2. The hydrogen isotopes were observed in a redeposited carbon layer.

5. Interaction with oxygen and influence on impurities in tokamaks The use of boron carbide materials has lead to substantial improvement of plasma parameters when compared to materials with carbon only mainly because of reduction of oxygen impurity by means of oxygen gettering [33]. It was shown that oxidation of boron carbide coatings took place under the influence of energetic oxygen ions even at room temperature [5,34]. This process has been investigated in some detail. Polycrystalline B4C prepared by sintering B4C powder through hot pressing and 2.5 keV oxygen ions were used in ref. [35], and plasmasprayed B4C layers on graphite and 1 – 5 keV ions were used in ref. [36]. At the beginning of a steady state irradiation, almost 100% oxygen trapping takes place. Then after a transient regime a steady state is reached where there is no oxygen gettering. CO and, in lesser amount, CO2 are reemitted

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during irradiation. Boron oxides sputtering is also observed but their contribution amounts to only a few percent [36]. The retained oxygen is contained in the form of nonvolatile boron oxides BO, BO2, B2O2, B2O3 [35,36], probably as B–C–O complexes [36] and also physically trapped CO and O2 [35]. Interaction of B4C with water is also of importance. Remarkable oxidation was observed under simultaneous irradiation by energetic electron and helium ion beam, but no oxidation took place under the same irradiation conditions when O2 was used instead of H2O [35]. Interaction of B4C preliminary irradiated by deuterium ions with water has also been investigated [37]. Porosity is developed under irradiation which is responsible for water vapor penetration into the bulk material and formation of B2O3 as a result of interaction of H2O and B4C at a room temperature. An oxygen concentration in JT-60U gas converted isotropic graphite tiles has also been studied [26]. The highest concentration was observed on the surface of the first wall outboard upper tile. It was more than two times higher than in graphite tiles, i.e. B4C layer was more effective as an oxygen getter compared with carbon. The performance of gas converted CFC tiles has been investigated in high power neutral beam heated divertor discharges of JT-60U with a heating power of  30 MW for 2 s [14]. 100 mm thick layers with a B–C ratio of 2.7 reduced carbon impurity by 3 times without a significant increase of boron, and oxygen impurity by 2 times. These improvements were in good agreement with the results obtained by boronization. Drastic reduction of the oxygen impurity was achieved by using 300 mm thick coatings with a B–C ratio of 3.8, though the significant melting and influx of boron took place. It was explained by solid target boronization due to evaporation of the tile. As a result of more pure plasma due to a fluoride CVD coating on the graphite rf-antenna, the effectiveness of ion cyclotron resonance heating increased by 3 times from 5 eV/kW to 15 eV/kW in the T-11 M tokamak [38].

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6. High heat flux tests Coatings should endure high heat fluxes without large erosion or damage, and exhibit superior surface characteristics for a sufficiently long time. There are two different heating conditions: the heat load of normal tokamak operation with a power density of some tens MW/m2 during several seconds and that of the thermal quench phase of a disruption with an energy density of some tens MJ/m2 in a millisecond or submillisecond time scale. Usually high energy electron, ion and laser beams are used to simulate such conditions. However it should be noted that experiments in plasma accelerators give an order of magnitude less erosion than laser or electron beam which is connected with shielding of a sample surface by a cloud of vaporized substance [39]. Due to high energy ( 100 keV) and relatively large mean free path of electrons in low Z materials the shielding effect can not be simulated in electron beam facilities. Besides, in this case a part of incident energy is absorbed by the underlying substrate. More correct experiments were performed on plasma beam devices PLADIS (15 MJ/m2, 0.1 ms) and 2MK-200 (1.5 MJ/m2, 0.02 ms) with fluoride CVD coatings [13] although the parameters of the latter device seems somewhat low. The results of the high heat flux tests for coatings of all types are summarized in Tables 1–3.

6.1. Chloride CVD coatings The performance under normal operation heat loads produced by a hydrogen ion beam was investigated at KFA [8]. The pulse duration was increased from shot to shot for a given power density until a failure of a specimen (melting, cracking or severe erosion) occurred. When the power density and the pulse duration were 6 MW/m2 and 6 s, respectively, the 0.2 mm thick coatings developed a typical crack pattern due to differences in thermal expansion coefficients of the substrate (C – C composite CC1001G) and the coating. The adhesion remained good. Experiments on the JEBIS stand with heat fluxes up to 24 MW/m2 and pulse duration 5 s have been carried out [9]. Contrary to above

mentioned studies coatings on composite CC312 did not show any damages up to 12 MW/m2 corresponding to normal conditions of the JT-60 divertor. Melting, cracking and color changing of the B4C layer occurred at fluxes over 15 MW/m2 so that up to 20 MW/m2 there was only color changing. Although exfoliation did not occur at applied heat fluxes, it was supposed that it might occur under higher power densities. X-ray diffraction analysis did not revealed any changes in the B–C ratio because it averages data over a large depth. Electron probe microanalysis capable of analyzing an atomic composition in a surface layer of some micrometers in depth showed that the B–C ratio substantially (by 1.5 to 2 times) decreased due to preferential vaporization of the boron atoms at high (\ 800°C) temperatures. Disruption simulations were performed on the JEBIS stand with a power density of 550 MW/m2 and a pulse duration of 5–10 ms (2.7− 5.5 MJ/ m2) [9]. The chloride CVD coatings were severely exfoliated at the interface between the layer and the substrate.

6.2. Fluoride CVD coatings The behavior under normal heat loads was investigated under electron beam irradiation at Sandia National Laboratory [13]. The coatings on the MPG-8 graphite had no noticeable damages up to 10 MW/m2 when erosion occurred. Profilometry showed local damages with maximum crater depth of 180 mm and the thickness of the remained layer 200 mm. The depth of the other small defects was about 30 mm. More than half of the surface was intact. The coating on the RGT graphite was preserved up to the maximum applied power density of 13 MW/m2 due to the high thermal conductivity of this graphite and hence lower surface temperature as compared to the MPG-8 coatings. The X-ray analysis revealed no changes in phase content, not even in the bottom of the largest crater, although some boron enrichment was found in the irradiated surfaces. The fluoride CVD coatings have been also investigated in tokamaks: as a limiter in T-10 and divertor tiles in DIII-D. In T-10, the discharges were both ohmic and with electron cyclotron reso-

Table 1 Heat load tests: normal conditions Thickness (mm)

Chloride CVD

CC1001G

200

Chloride CVD

CC312

100, 250

Fluoride CVD

MPG-8

400

Fluoride CVD

RGT

100

Gas conversion PCC-2S

50–200

Gas conversion CFC composite Plasma spray CC1001G

80, 140 200

Plasma spray

200

Plasma spray Plasma spray Plasma spray Reaction-sintering

EK98, AXF 5Q, KKARB PCC-2S, CX2002U, MFC-1 Copper based AISI316L with Cu interlayer EK98

150 Up to 200 1200 50–1000

Device/type of heat flux

Damage threshold

Damage type

Ref.

20–50 keV hydrogen ion beam JEBIS/100 keV electron beam Sandia Nat. Lab/Electron beam Sandia Nat Lab/Electron beam JEBIS/100 keV electron beam 20 keV electron beam 20–50 keV hydrogen ion beam 20–50 keV hydrogen ion beam JEBIS/100 keV electron beam FE-200/150 keV electron beam JEBIS/Electron beam

6 MW/m2,6 s

Crack

[8]

20 MW/m2, 5 s

[9]

10 MW/m2, 2 s

Crack, exfoliation Erosion

[13]

No up to 13 MW/m2, 5s

–

[13]

15 MW/m2, 5 s

Melting

[9]

20 MW/m2, 60 s 6 MW/m2, 6 s

Melting, crack Crack

[15] [8]

6 MW/m2, 6 s

Melting

[8]

15 MW/m2, 5 s

Crack, exfoliation –

[9]

[23]

JUDITH/140 keV electron beam

12 MW/m2, 3 s

Crack, delamination Melting

No up to 7 MW/m2, 30s (15 MW/m2, 3 s for CuCrZr) 0.7 MW/m2

[24]

[19]

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Method of fab- Substrate rication

351

352

Methods of fab- Substrate rication Fluoride CVD

Thickness (mm)

Fluoride CVD Fluoride CVD Fluoride CVD Gas conversion

1)MPG-8, 2)Pyrolytic graphite MGP-8 RGT TS-1792 CFC composite

80–100

180 180 85–180 100–300

Plasma spray

EK98

100

Reaction-sintering

EK98

50–1000

Tokamak/element

Power density/pulse duration

Quantity of discharges/time of exposure

Damage

Ref.

T-10/limiter

5 MW/m2, 0.5–1 s (ohmic), 20 MW/m2, 0.1–0.4 s (ECR)

2500 (including some hundred with disruption)

1)Melting 2)Erosion

[11,12]

DIII-D/divertor DIII-D/divertor DIII-D/divertor JT-60U/Divertor TEXTOR/limiter TEXTOR/limiter

0.4 MW/m2 10–12 MW/m2 – 8–17 MW/m2, 1s

11 14/11 s – 1000

No Erosion Minor erosion Melting

[13] [31,32] [29] [14]

Low

11

Melting

[18]

12–26 MW/m2

15

Melting

[18,19]

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Table 2 Heat load tests: real tokamak conditions

Table 3 Heat load tests: disruption conditions Thickness (mm)

Device/type of heat flux

Energy density/pulse duration

Damage type

Ref.

Chloride CVD

CC312

100, 250

[9]

MPG-8 RGT MPG-8

100 100 100

550 MW/m2, 5–10ms (2.7–5.5 MJ/m2) 15 MJ/m2, 0.1 ms 15 MJ/m2, 0.1 ms 1.5 MJ/m2, 0.02

Exfoliation

Fluoride CVD Fluoride CVD Fluoride CVD

Erosion, melting, cracks Erosion, microcracks Minor erosion

[13] [13] [13]

Gas conversion

PCC-2S

[9]

Exfoliation

[8]

Plasma spray

EK98, AXF 5Q, K-KARB, CC1001G, steel 1.4311, molibdenum PCC-2S, CX2002U, MFC-1

550 MW/m2, 5–10 ms (2.7–5.5 MJ/m2) 6 MJ/m2, 5ms

Melting

Plasma spray

JEBIS/100 keV electron beam PLADICE/plasma beam PLADICE/plasma beam 2MK-200/1 keV plasma beam JEBIS/100 keV electron beam JEBIS/100 keV electron beam

150

[9]

EK98

100

550 MW/m2, 5–10 ms (2.7–5.5 MJ/m2) 2.5–6.2 MJ/m2, 2 ms

Exfoliation

Plasma spray Plasma spray

PCC-2S, CC1501GR, 316L

240–800

Plasma spray

316L

1200–1500

Plasma spray

316L

1200–1500

Reaction sintering

EK98

JEBIS/100 keV electron beam JUDITH/140 keV electron beam JEBIS/80 keV (150 kW) electron beam JUDITH/100 keV (B24 kW) electron beam FE-200/100–160 keV (40–52 kW) electron beam JUDITH/140 keV electron beam

50–200 200

50–1000

Erosion, crack, melting, damage 2.15 MJ/m2, 1.2 ms Craters, chipping, damage of thinnest 1.8–2.4 MJ/m2, 1&10 ms Melting, chipping, no complete failure 6–9 MJ/m2, 1–10 ms Craters, chipping Up to 8 MJ/m2, 2 ms

Melting

[19] [20–22] [21,22] [21,22] [18,19]

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Methods of fab- Substrate rication

353

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Fig. 8. The thicknesses of eroded layers of limiter coated with various materials.

nance (ECR) heating, and some hundred of them was with disruptions [11,12]. Mass losses of the samples with B4C coatings were 1.5 – 2 times lower than those of the USB-15 graphite (Fig. 8). The typical initial thickness of the B4C layer was 80– 100 mm and post exposure thickness was 10–30 mm. Melting of the surface layer followed by recrystallization and the formation of a crystallite grid was observed on the surface of the coated MPG-8, especially on the electron side of the limiter. It means that the surface temperature was very high (\ 2000°C). Significant melting of the B4C coating on pyrolytic graphite was not observed though the heat and particle fluxes on this coating were substantially higher. It can be explained by the high thermal conductivity of the pyrolytic graphite (about 300 W/m K) as compared to that of MPG-8 (70 – 80 W/m K). The coatings on all substrates had no macrocracks typical of the exposed graphite. Both initial coatings and exposed ones had good adherence to the graphite substrates. X-ray diffraction analysis and Auger spectroscopy showed that the phase and chemical composition of the coatings were preserved: the exposed and initial ones were mixtures of B4C and amorphous carbon. Some carbon enrichment in the surface layers was observed which can be explained by thermal desorption of boron at high temperatures. The DIII-D experiments were carried out in accordance with the DIMES (divertor materials evaluation system) program [40]. The coatings were exposed to the divertor strike point plasma flux. The coating on the MPG-8 graphite was

exposed to 11 discharges with low enough heat fluxes of about 0.4 MW/m2 [13]. The essential erosion was not observed and the surface topography after the exposure did not changed almost. The coating on the RGT graphite was exposed to 14 discharges [31,32]. The accumulated exposure time was about 11 s of ELMing H-mode plasma with some additional energy deposition from a locked mode disruption that terminate one of the discharges. The average heat flux at the sample radius measured by an infrared camera was 2–3 MW/m2. However because of not good alignment of the sample the peak flux was estimated to be about 10–12 MW/m2, somewhat above those loads expected for ITER. A damaged area appeared in the region of the highest heat flux. The thickness of the remaining B4C layer was about 130 mm and by comparison with neighboring undamaged layer about 45 mm was eroded. Any changes in crystalline structure, phase and chemical composition were not observed. Divertor tiles from the B4C-coated TS-1792 graphite were exposed to the DIII-D plasma during 16 operation weeks of the 1993 year’s campaign [29]. The thickness of the coatings varied from 85 to 180 mm. Typically, the plasma current was 0.5–2.5 MA, pulse duration was about 5 s and the power of neutral beam heating varied from 2 to 16 MW. The crystalline structure, phase and chemical composition of the coatings did not change. There were observed some minor defects such as craters and holes, created by the plasma fluxes, arcs and thermal shocks. The thickness of the eroded layer varied from 40 to 70 mm. In the largest defect, the thickness of the remaining coating was 50 mm. It should be noted that the tile surfaces were overcoated with a redeposited layer consisting mainly of carbon. Disruption simulations were performed on the PLADICE plasma beam device at the University of New Mexico [13]. Profilometry of the coating on the MPG-8 graphite showed that 80 of 100 mm was eroded. The phase content did not change, but some boron enrichment was detected. There were observed some melted areas and cracks. The coatings on the RGT graphite were less damaged showing some cracks and holes. The depth of the largest defect was 64 mm and everywhere the

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substrate was not reached. Scanning electron microscopy showed some cracks healed by melting and crystallization. Any changes in the phase content were not observed. Disruption simulations were also performed on the 2MK-200 plasma accelerator at TRINITI in Russia [13]. A coating on the MPG-8 graphite was exposed to 22 pulses with a total energy of 33 MJ/m2. No considerable defects were observed. The thickness of the eroded layer was about 10 mm, and there were some melted areas. The phase contents was not changed.

6.3. Gas con6ersion coatings Normal operation behavior of converted PCC2S composite with B/C =2.9 was investigated on the JEBIS stand [9]. As well as for the chloride CVD coating, no failure was found at fluxes up to 12 MW/m2. Melting occurred at power densities more than 15 MW/m2. Similar to the chloride CVD coating, there was observed substantial reduction of boron content in the surface layer of several micrometer depth but no exfoliation was observed up to 30 MW/m2 contrary to supposed behavior of chloride CVD coating. Two type of converted CFC composite samples with B/C equal to 2.7 and 3.9 similar to the JT-60 divertor tiles were investigated under electron beam with power density 20 MW/m2 [15]. Weak melting of the edge was observed at the surface of the first specimen. It was explained by locally increased surface temperature due to low thermal diffusion. The central region of the second specimen was melted and resolidified. Cracks were produced on it because of thermal stress after resolidification. It should be noted that no modification occurred on the surface of nonconverted substrate under such conditions because of its lower temperature due to higher thermal conductivity. The performance of the converted CFC composite divertor tiles were investigated in high power neutral beam heated divertor discharges of JT-60U with a heating power 30 MW for 2 s [14]. On the outboard strike point 100 mm thick coatings with B/C =2.7 and on the adjacent row 300 mm thick coatings with B/C =3.8 were in-

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stalled. Slight erosion was observed on almost all the tapered edges of the tiles after  1000 shots. Under an edge heat flux of 8–17 MW/m2 and a pulse duration of 1 s the edge temperature was estimated to be 1300–2400°C, and radiation enhanced sublimation might take place under these conditions. Significant melting of the coatings surfaces and influx of boron were observed. It was found that arching occurred frequently at the termination of the discharge, and melting of the layer was probably triggered by this arching. In addition it was noted that the resolidified melted layer forms a rugged surface probably due to effect of surface tension. This roughness enhanced the heat flux received by the surface because of the shallow angle of particle incidence. This was probably the reason why melting and splashing of the layer occurred and rapid boron influx was observed. However after a few hundred shots repetitive melting and splashing of the layer flattened the rugged surface and boron influx did not occur. Disruption simulation were performed on the JEBIS stand under conditions similar to that for chloride CVD coatings but contrary to them no exfoliation was observed although the surface was partially melted [9].

6.4. Plasma spray These coatings have been extensively investigated. The performance of LPPS layers under normal heat loads was similar to that of chloride CVD layers in ref. [8]. The 0.2 mm thick coatings on the CC 1001 G composite developed a typical crack pattern due to differences in thermal expansion coefficients between them and the substrates but adhesion remained good. When the matching of the expansion coefficients was much better the coatings did not show any cracking and heat load could be increased up to the melting point of B4C. Normal operation behavior of LPPS coatings with B/C = 3.6 was investigated on the JEBIS stand under heat fluxes up to about 40 MW/m2 and pulse duration 5 s [9]. Similar to chloride CVD and gas conversion coatings, no failure was found under heat fluxes up to 12 MW/m2. Differ-

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ences began at densities more than 15 MW/m2 when LPPS coatings began to crack and at 30 MW/m2 to exfoliate so that the coating on PCC2S exfoliated even earlier at 24 MW/m2. The weight losses of a sample was drastically increased when exfoliation occurred because of poor heat conductance and the layer burned through. There was observed substantial reduction of boron content in the surface layer of several micrometers in depth due to preferential vaporization of boron at high temperatures. Thermal cycling behavior of plasma spray samples and a full size actively water cooled TORE SUPRA port protection element were investigated in the electron beam facility FE-200 [24]. Some plasma facing components such as Faraday screen of ion cyclotron heating antenna, waveguide mouth of low hybrid wave antenna, port protections, divertor plates have complex geometry so that their shaping and CFC brazing is difficult. Therefore it was supposed to produce them from copper based materials and then to cover them with boron carbide by the plasma spray technique. All the coatings were deposited without gradient layers. Their thermal conductivity was about 1 W/m K which is low compared to the bulk boron carbide (30 – 40 W/m K).The thickness of the coating on the port protection element was from 30 mm on edges to 60 mm on flat parts. After 1000 cycles at 7 MW/m2 (30 s on, 30 s off) no defects were observed. Due to relatively large mean free path of electrons, the energy deposited in the layer was evaluated to be only 25%. Coating on actively cooled DS copper substrate were damaged after 800 cycles (3 s on, 3 s off) at 15 MW/m2. Similar experiments with a coating on CuCrZr substrate showed good thermal cycling behavior due to a better bonding quality. However, coatings thicker than 200 mm on copper based substrates without a gradient layer had adherence problems. It was stated that the use of such relatively thin coatings might be acceptable for such tokamak as TORE SUPRA because replacement of components is not a major problem. However, for future large machines where thicker coatings are required thermal conductivity must be substantially enhanced.

Poor thermomechanical properties due to edge stresses between the layer and the substrate were observed in ref. [23]. As substrates, actively cooled stainless steel (AISI316L) specimens were used which corresponded in shape to a simplified ITER-CDA design of the first wall. An approximately 1.2 mm thick B4C layer was deposited on top of the Cu interlayer with a thickness of about 3 mm. The electron beam irradiation of the specimens was carried out in the JEBIS facility. One of the two identical specimens was exposed to a sequence of pulses with increasing power density and with varying duration. The other was subjected to pulses of defined power density (0.5 MW/m2) and duration (150 s). During thermal loading both specimens showed one or two hot spots at the edges which were visible at the lowest heat fluxes. On the first specimen the intensity of the hot spots increased with increasing flux. At 0.7 MW/ m2 the coating cracked at the hot spot and progressive delamination of the B4C coating from the Cu interlayer took place during subsequent pulses at 1 MW/m2. At the edges of the second specimen hot spots formed progressively with increasing pulse numbers. During the tenth pulse delamination of the coating started at the hottest spot. These spots at the free edges of the specimens indicate to preexisting delamination which appears to be a consequence of the stress singularity formed at the edge during cool down at the end of coating deposition process. Since the thermal expansion coefficients of Cu and B4C are strongly different, the metallic components contract more than boron carbide during the initial cool down and causes the localized debonding. Under surface heat loads the less well cooled B4C in the locations of initial delamination is subjected to a much higher thermal excursion than the cooled Cu substrate below the delaminated zone. The resulting stronger expansion of the B4C leads to compressive stress in it and coating tends to flex off from the interlayer. It was stated that geometrical modifications should be undertaken to decrease the influence of the stress singularity at the free edge of the coating–interlayer–substrate interfaces, and the interface itself should be geometrically optimized to improve the adhesive strength.

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A porous almost stoichoimetric 100 mm thick coating on the EK98 graphite was used as test limiter in the TEXTOR tokamak [18]. Typical discharge duration was 3.5 s, they were either ohmic or with neutral beam injection of up to 1.2 MW. During subsequent discharges the test limiter was moved deeper into the plasma. The total number of exposures was 11. Energy flux onto the limiter surface was calculated from the comparison of its temperature in plasma and under electron beam irradiation. However surface temperature measurements were not available for LPPS coating, and it was only stated that the coating melted under low heat flux due to its low thermal conductivity evaluated to be about 4 W/ m K. The scanning microscopy and electron diffraction analysis showed that even though the melted coating had agglomerated to droplets on graphite surface, the whole surface area was infiltrated by liquid B4C. The disruption simulations were performed in the JEBIS [8,9,20 – 22], JUDITH [19,21,22] and FE-200 [21,22] electron beam facilities, and with Nd: YAG laser beam [21,22]. The 200 mm thick LPPS coatings had a very limited capability to withstand disruptions in the MJ/m2 range in ref. [8]. It was observed their partial removal and melting of the substrates. Similar results were obtained in ref. [9] where the 150 mm LPPS coatings were severe exfoliated at the interface between the layer and the substrate. In ref. [19], an erosion crater and microcracks formed at 2.5 MJ/m2, and melting beginning at relatively low energy densities caused the damage of the LPPS coatings at 6.2 MJ/m2. More thick coatings were investigated in [20– 22]. The thickness of coatings sprayed directly on the stainless steel substrate was limited to 0.3 mm because the strong difference in thermal expansion coefficients induced stresses and detachment. The interlayer with gradually decreased steel fraction about 0.4 – 0.5 mm in thickness were sprayed onto the substrate. The B4C top coatings up to 1.6 mm were subsequently sprayed. Significantly higher weight losses were measured on coatings on stainless steel when compared to CFC due to its poor thermal conductivity. The absolute mass loss was much

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higher than that of the uncoated carbon materials, especially when the low energy densities values corresponding to disruption heat loads on the first wall were taken into account. Similar to ref. [8,9], the thinnest 240 mm coating was damaged to underlying substrate. Contrary to carbon materials, all B4C coated specimens showed large erosion craters with diameters up to 11.3 mm (beam diameter 5.3 mm), i.e. significant erosion also occurs in areas with relatively low energy deposition on the edge of the beam. The typical crater depth of the coatings on 316L was about 0.12– 0.18 mm maximum and 0.05 mm on average. The crater profile revealed a stepped structure with steps of about 0.05–0.06 mm connected with that the coatings were produced by spraying subsequent layers. So there was observed chipping, i.e. exfoliation of the individual layers of the coating rather than sublimation of the coating material. The thresholds for observable damage was found by stepwise increase of power in ref. [21,22]. They were about 0.9 MJ/m2 for 1 ms and 2.1 MJ/m2 for 10 ms. Typical surface effects were sintering, melting and chipping. Particle emission were visible on the monitoring video camera. Some samples which performed well in the initial tests were loaded with 80–100 shots at 1.8 and 2.4 MJ/m2 for 10 ms. After each shot they were given time to cool down to about 250°C. Although significant surface damage (melting, chipping) occurred, no complete failure of the coating was observed. The total crater depth was in the range of 0.3 to 0.5 mm. In the FE-200 tests [21,22], the resulting crater on the coating with 0.5 mm interlayer and 1.2 mm top layer was 0.3 mm maximum in depth and showed the stepped structure similar to JEBIS tests. Experiments with thinner top coating showed craters of 0.1 mm depth for 2.1 MJ/m2 and 0.2 mm for three shots at 3 MJ/m2. Experiments with Nd:YAG laser [21,22] showed that the thresholds for surface melting (about 0.35 MJ/m2 for 1 ms pulse and 1.3 for 10 ms pulse) were lower than those under electron beam irradiation which can be mainly attributed to the effects of volumetric energy deposition of electrons.

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6.5. Reaction-sintering The coating evaluation tests included plasma heating and disruption simulation experiments in JUDITH and plasma heating experiments with a coated test limiter in TEXTOR. In order to simulate the plasma start up and ramp down phases of tokamak discharges, the electron beam power was ramped up and down linearly for 0.5 s with flat top phase 3 s [19]. The heat flux was increased in steps. During heating at 12 MW/m2 melting of the coating commenced at the end of the 3 s flat top condition. Heating at the 15 MW/m2 level caused melting of the coating after 1.8 s. After exposure, the beam heated zone was characterized as a flat droplet covering about half of the area. The other half was covered with a thinner coating. The thermal conductivity was higher than 4 W/m K mentioned above for LPPS coatings. The B4C coated limiter was exposed to 15 plasma discharges in TEXTOR [18,19]. Initial melting occurred at power density 12 MW/m2. At the highest loading of 26 MW/m2 melting took place after 0.5 s. Temporal behavior of the surface temperature under high heat loads (when melting occurred) in tokamaks differed from that under electron beam irradiation. In the TEXTOR tests as well as in the electron beam tests the rise of the surface temperature was intermitted when surface began to melt. In the case of the electron beam this can be partly ascribed to the formation of the melt pool. In TEXTOR the melted zone was much less, i.e. melting is of minor importance with regard to energy balance. In this case the evaporated species may exert a shielding effect which causes a reduction of the incident heat flux. After resolidification carbon was segregated in the form of needles and small platelets at the B4C grain boundaries. During exposure the test limiter fractured into fragments due to mechanical and thermomechanical stress. Under the disruption simulations, penetration depth of electrons was 130 mm [19]. Under the highest heat load conditions (\6.2 MJ/m2) the B4C–C coatings were more resistent to erosion

and melting than an uncoated finegrain graphite and LPPS coatings heated less severely. No significant coating degradation was observed. The coatings exhibited eutectic melting near 2400°C under rapid heating. Adherence was maintained with a reduced release of carbon at high temperatures.

7. Conclusion Plasma contamination with impurities and protection of plasma facing components (PFC) against high heat fluxes are one of the most important issues in present-day tokamaks. To solve these problems several methods can be applied. One of them is protection of PFC using boron–carbide coatings. Contrary to pure carbon materials, B4C has much lower chemical and high-temperature sputtering and possesses such important features as oxygen gettering and lower hydrogen recycling. In contrast to thin boronization films the thick coatings are capable to resist high heat fluxes such as in tokamak divertors. There are several methods to obtain such coatings: chemical vapor deposition by means of chloride and fluoride techniques, gas conversion, plasma spray and reaction-sintering. The main demand on protection elements is the ability to withstand high heat loads from the main plasma. Usually high energy electron, ion and laser beams are used to simulate such conditions. However experiments in plasma accelerators give an order of magnitude less erosion because of the shielding of a sample surface by a cloud of vaporized substance. More correct experiments have been performed in plasma beam devices and directly in tokamaks. Comparison of all available data shows that coatings produced by diffusion methods, such as fluoride CVD and gas conversion, are more resistent to heat loads, and one of the most promising candidates are the fluoride CVD coatings. They have high density, crystalline structure, good adhesion, preserve their properties and protect plasma facing components against high heat plasma fluxes.

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