**3. Carbide dispersion-strengthened tungsten-based materials**

As compared with the oxides' strengthening phases, the carbides, such as TiC, ZrC, and HfC, have much higher melting temperatures and better compatibility with tungsten, which may lead to excellent comprehensive performance. Kurishita et al. reported ultrafine-grained (UFG) W-TiC alloys which showed increased

mechanical properties and better irradiation resistance [22, 39, 40]. These favorable properties of W-TiC alloys imply the potential of other carbide dispersion-strengthened tungsten materials with fine microstructures.

Among these carbide-strengthened phases, ZrC and TaC have high melting temperatures of 3540°C and 3900°C, respectively. It's worth mentioning that the lattice match of *d*(200)ZrC ≈ *d*(110)W ≈ 0.221 nm might introduce coherent PB interface between ZrC phase and W matrix, which will significantly increase PB cohesion. In addition, ZrC as an oxygen getter can react with oxygen to form stable Zr-C-O or ZrO2 particles at GBs, purifying GB interface, and thus is beneficial to improve the GB cohesion and enhance the low-temperature ductility of tungsten.

Fan et al. [27, 41] fabricated W-ZrC material by high-temperature sintering in hydrogen atmosphere using W-ZrC composite powders prepared by sol-heterogeneous precipitation-spray drying-thermal reduction. The relative density of their W-(0, 1, 2, 3, 4)%ZrC samples is in the range of 98.5–99.7%. They found that the addition of ZrC could refine tungsten grains and improve the strength. The grain size of W-3%ZrC (10~15 μm) is much smaller than that of pure W (~100 μm), and the ZrC particles are micron-sized; most of them are distributed at grain boundaries. Xie et al. [42, 43] fabricated a series of W-(0, 0.2, 0.5, 1.0)%wtZrC with an average grain size which ranges from 2.7 to 4.2 μm by using mechanical alloy and SPS. The ultimate tensile strengths (UTS) of SPS pure W and W-(0.2, 0.5, 1.0)%ZrC at 700°C are 337, 419, 535 and 749 MPa, respectively, suggesting enhancements in strength by adding a small amount of ZrC nanoparticles. The strength of W-0.5%ZrC (W-0.5ZrC) is over 50% higher than that of pure W. Meanwhile, the ductility could also be effectively improved by the ZrC addition.

On the basis of composition optimization, Xie et al. [28] fabricated bulk W-0.5wt%ZrC plates with a thickness of 8.5 mm, which are suitable for engineering application. This material exhibits very good mechanical properties. At RT, this bulk W-0.5ZrC plate exhibits a high bending stress of 2.5 GPa as well as a flexural strain of 3%, which is much higher than those of the hot rolled W, HIPed pure W [44], and W-1.0Y2O3 [45] and close to those of the severely deformed W-TiC alloys [46]. The DBTT is about 100°C (see **Figure 1** and **Table 1**), which is much lower than that of reported bulk W alloys and several other tungsten materials [31, 47]. At temperatures above 150°C, the W-0.5ZrC plates can be bent to a high flexural strain of 15% (limited by the machine) without any crack.

**Figure 1e** shows the engineering stress-stain curves of the W-0.5ZrC alloy tested at various temperatures along rolling direction (RD). At 100°C, the W-0.5ZrC plate exhibits obvious tensile deformation with a TE~3% and an UTS value up to 1.1 GPa. With the increase of the test temperature to 200°C, the TE increases to 14.2%, and the UTS maintains as high as ~1 GPa. At a higher testing temperature of 500°C, the UTS is still very high (583 MPa), and the TE increases up to 41%. The rolled W-0.5ZrC showed both higher strength and ductility than that of the rolled W-0.5TiC [48] and rolled W-0.5TaC [49], which were fabricated following a similar process. The W-3%Re and K-doped W-3%Re show ductility at 100°C, which is comparable with W-0.5ZrC plate, while their tensile strength values (<1000 MPa) are lower than that of W-0.5ZrC plate (1058 MPa) [50]. As compared with Plansee ITER specification W (IGP), AT&M ITER specification W (CEFTR), W-1wt%TiC (W1TiC), W-2wt%Y2O3 (W2YO) from Karlsruhe Institute of Technology at Germany, and fine-grained W (FG) from the Institute of Plasma Physics at Czech Republic, W-0.5ZrC exhibits the highest tensile strength and lower DBTT [51], as shown in **Figure 2**.

For PFCs, the transient heat load events such as major plasma disruptions, edge-localized modes (ELMs), and vertical displacement events (VDEs) are serious issues [11], which may lead to a significant temperature rise and high thermal

*The Tungsten-Based Plasma-Facing Materials DOI: http://dx.doi.org/10.5772/intechopen.88029*

#### **Figure 1.**

*Mechanical properties of W-0.5ZrC alloy: (a) flexural stress-strain curves tested at various temperatures (note that values larger than a flexural strain of 15% is not accurate due to the limited bending angle of the machine). (b) Flexural strain of W-0.5ZrC at various test temperatures. (c) Temperature dependence of the yield strength (YS) by 3-point bending test in comparison with available literature data. (d) Optical images of 3-point bending specimens tested at various temperatures. (e) Tensile engineering stress-strain curves of W-0.5ZrC at different temperatures. (f) Hardness in different planes of W-0.5ZrC tested by nanoindenter [28].*

stresses in PFCs. The abrupt high temperature during transient events can lead to material recrystallization, grain growth, surface melting, and droplet ejection. The high thermal stresses can lead to cracking, fatigue fracture, and fatal destruction of the PFCs. Therefore, the thermal shock resistance of PFCs is closely related to their mechanical properties. Intuitively, high strength could resist a relatively high stress induced by thermal shocks to prohibit the cracking formation, while good plasticity/ductility is in favor of releasing the stress via plastic deformation rather than cracking. That is to say, a higher strength and better low-temperature ductility would lead to the better thermal shock resistance [28, 31, 52].


#### **Table 1.**

*DBTT of the rolled W-0.5ZrC and several reported W materials [31].*

#### **Figure 2.**

*The tensile stress of various W materials [51].*

**Figure 3** shows the thermally loaded surfaces of rolled W-0.5ZrC after exposure to single shot with a pulse length of 5 ms using an electron beam. No cracks or melting were observed on the samples when tested at an absorbed power

*The Tungsten-Based Plasma-Facing Materials DOI: http://dx.doi.org/10.5772/intechopen.88029*

#### **Figure 3.**

*The SEM images of sample surfaces after exposure to thermal shocks with a pulse duration of 5 ms. (a) 3.3 MJ/ m2 (0.66 GW/m2 ), (b) 4.4 MJ/m2 (0.88 GW/m2 ), and (c) 5.5 MJ/m2 (1.1 GW/m2 ) [28]. No cracks were found at 0.88 GW/m<sup>2</sup> although the sample surface melted and the cracking threshold is 0.88~1.1 GW/m2 .*

density (APD) of 0.66 GW/m<sup>2</sup> as shown in **Figure 3a**. Still no crack at an APD of 0.88 GW/m<sup>2</sup> but surface melting was observed as shown in **Figure 3b**. The results suggest that the melting threshold of the rolled W-0.5ZrC is ~0.88 GW/m<sup>2</sup> . When the APD increased to 1.1 GW/m<sup>2</sup> (**Figure 3c**), both cracks and melting were found on the sample surface, indicating a crack threshold of 0.88–1.1 GW/m<sup>2</sup> . For comparison, the cracking thresholds of previously reported tungsten materials like sintered W, chemical vapor deposition (CVD) W, deformed W, potassiumdoped W (W-K), and W-La2O3 alloys [53–57] were summarized in **Table 2**. These results further indicate the excellent thermal shock resistance of the rolled W-0.5ZrC.

During the thermal shock, the materials are heated and undergo thermal expansion, which is restricted by the colder surrounding material, causing compressive stress [53]. If the compressive stress exceeds the yield strength of the tested tungsten material, plastic deformation would occur because of the material plastic at high temperatures. During the cooling stage, the thermally loaded area shrinks, and the compressive stress is converted into a tensile stress [53]. Thus the cracks would form if the stress exceeds the ultimate tensile strength of the materials. Therefore, the excellent thermal shock resistance of the rolled W-0.5ZrC seems reasonable, because it has high strength, a lower DBTT, and good plasticity.

Thermal fatigue behaviors of the as-rolled and recrystallized W-0.5ZrC alloys were investigated by repetitive thermal shocks (100 shots in total) with a pulse duration of 1 ms at RT [14]. The cracking thresholds for the as-rolled and recrystallized W-0.5ZrC are 0.22–0.33 GW/m2 , which is comparable to that of ultrahighpurity W but lower than that of W-5%Ta [52]. It's reported that at higher base temperatures, tungsten materials tend to exhibit enhanced resistance to thermal


#### **Table 2.**

*Comparison of W-ZrC alloys and some tungsten materials after single thermal shock loads at RT with a pulse duration of 5 ms [31].*

shocks because of the increased ductility of materials [13]. Nevertheless, there are still no results of the thermal fatigue behavior of the rolled W-0.5ZrC at higher base temperatures, which would be investigated in the near future.

The irradiation resistance to plasma is another key property for PFCs. The PFCs in fusion reactors suffer from the high-flux (about 1024 m−<sup>2</sup> s<sup>−</sup><sup>1</sup> ) low-energy plasma irradiation, which would cause bubbles and erosion phenomena on the surface of PFCs, and lead to the degradation of performances. More badly, the excessive erosion dusts will extinguish the burning plasma. Therefore, the less the erosion, the better.

Liu et al. [58] studied the irradiation damage of several newly developed tungsten materials including pure tungsten, CVD-W, W-0.5ZrC, W-1.0wt.%Y2O3, W-1.0vol.%Y2O3, and W-1.0La2O3 under low-energy He plasma neutral beam. **Figure 4** shows the surface and cross-sectional morphologies of irradiation-modified layers on the abovementioned tungsten materials after high fluences (1026 ions/m2 ) of low-energy helium plasma irradiation [58]. After 220 eV He-ion irradiation at 900°C, all samples showed pinhole features on the irradiated surface, while in the case of 620 eV He-ion irradiation at 1000°C, the pinhole surface evolved into corallike features except for the W-0.5ZrC alloy which retains the pinhole feature [58]. This result indicates the good resistance of W-0.5ZrC alloy to He-ion irradiation. The fine-grained W-0.5ZrC alloy has abundant GBs and PBs which all provide nucleation sites for He atom aggregation and thus reduce the concentration of He, thus hindering the growth of He bubbles and mitigating the evolution from pinhole to coral-like structures. The thickness of modified layers on the surface of tungsten materials was measured and plotted in **Figure 5**, in which W-La, W-Y1, and W-Y2 represent W-1wt.%La2O3, W-1wt.%Y2O3, and W-1vol.%Y2O3, respectively. The thickness of modified layers on W-0.5ZrC is much lower than that of pure W, W-La, and W-Y1, implying its better resistance to plasma irradiation and erosion [58].

Liu et al. [58] also studied the evolution of morphology and thermal-mechanical properties of pure W, CVD-W, and W-0.5ZrC alloys after pure H beam and H/ He mixed beam irradiation using the neutral beam facility GLADIS (Max Planck Institute for Plasma Physics, Germany). It is found that just roughness occurred on all W material surfaces after H irradiation, while a mixture of 6% He resulted in pinhole structures, indicating the crucial factor of He irradiation for the surface modification. The unexposed and pre-irradiated CVD-W and W-0.5ZrC were then

#### **Figure 4.**

*The surface morphologies of (a) pure W, (b) CVD-W, (c) W-1.0%Y2O3 irradiated by 220 eV He<sup>+</sup> at about 900°C and (d) pure W, (e) CVD-W, and (f) W-0.5ZrC irradiated by 620 eV He+ at 1000°C to a same fluence of 1 × 1026 atoms/m2 [58]. The insets of (a)–(f) show the corresponding cross-sectional morphologies of irradiation-modified layer in each sample.*

loaded repeatedly by thermal shocks through an electron beam with a 1 ms pulse and 100 cycles. The cracking threshold of unexposed CVD-W is about ~0.22 GW/m2 , a little lower than that of W-0.5ZrC. Pre-irradiation by H only seems to have less effect on the critical cracking thresholds, while pre- irradiation by H/He mixed beam significantly reduces the cracking thresholds [58].

The effects of deuterium (D) plasma irradiation on the microstructure of W-ZrC were also investigated. Several CDS-W materials, including rolled W-0.5ZrC [28], W-0.5HfC [30], W-0.5TiC [48], as well as pure W, were subjected to D plasma irradiation in the linear plasma device of simulation for tokamak edge plasma (STEP) [59]. **Figure 6** shows the morphologies of these materials after exposing to D plasma irradiation under a same condition

#### **Figure 5.**

*Thickness of modified layers in various tungsten materials under the He+ irradiation of 220 eV at 900°C and 620 eV at 1000°C [58].*

#### **Figure 6.** *Surface morphologies of pure W, rolled W-0.5ZrC, W-0.5HfC, and W-0.5TiC irradiated by D plasma. [31].*

(D<sup>+</sup> energy ~90 eV, flux ~5 × 1021 ions/m2 s, fluence~7.02 × 1025 ions/m2 , temperature ~180°C) [31]. Surface swelling and a high density of large bubbles with size ranging from 1 to 10 μm were found on the surface of the pure W. On the surface of W-0.5HfC and W-0.5TiC, the bubble size was much smaller (less than 1 μm). In the case of W-0.5ZrC, only blisters (about 100 nm) were observed, despite the high blister density. These results further indicate the increased resistance to bubble formation to D plasma irradiation. The enhanced irradiation resistance may come from the fine grains and homogeneously dispersed nanoscaled particles in the rolled W-0.5ZrC, which provide a large number of GB and PB interfaces that act as effective sinks for irradiation-induced defects [60–62].

The hydrogen isotopes especially tritium retention behaviors in PFCs are also an important issue, which has an impact on the tritium fueling and the safety of fusion reactors and has not been completely understood. The hydrogen isotope retention behaviors of W-ZrC alloy were investigated through plasma irradiation and thermal desorption spectrum (TDS) analysis. The D plasma irradiation was carried out in the STEP device [59]. Rolled pure W and W-0.5ZrC were exposed to low-energy high-flux deuterium plasma under the same condition (D<sup>+</sup> energy~ 90 eV, flux ~5 × 1021 ions/m2 s, fluence~7.02 × 1025 ions/m2 , temperature 400 K). After irradiation, the TDS results have shown that the rolled W-0.5ZrC exhibits much lower hydrogen retention than that of pure W as shown in **Figure 7**. This result indicates that the hydrogen retention in W-ZrC is not increased although they contain the Zr element, because the Zr element exists in the form of stable carbide or oxide particles. Both the rolled W-0.5ZrC and pure W had a desorption peak at temperature around 560°C, and the peak intensity of W-0.5ZrC is much lower than that of pure W.

The above performance is closely related to the microstructure of the rolled W-0.5ZrC. The detailed microstructures show the coexistence of multi-scale interfaces in the rolled W-0.5ZrC plate. From the outer to inner space, in the first layer along the rolling direction, the average length of mother grains is about

#### **Figure 7.**

*TDS results of pure W and W-0.5ZrC after 90 eV high-flux (~5 × 1021 ions/m2 s) D plasma irradiation to a fluence of 7.02 × 1025 ions/m2 at 400 K [31].*

10 μm, and the width is about 1~3 μm, which is the micrometer scale GB interface [28]. Prolonged mother grains come from the multistep rolling deformation. In the second layer, there are equiaxed sub-grains in the matrix with average size of about 1 μm, which could be considered to the sub-micrometer scale GB interfaces. The sub-grains can be attributed to the deformation and dynamic recrystallization resulting from the precise control of rolling parameters [28]. For the third layer, most of the nanoscaled particles disperse in tungsten grain interior. The size distributions of the particles indicate that most particles located in W grains have an average size of 51 nm, while parts of particles at W GBs show bimodal distribution which contains relatively small particles with an average particle size of 60 nm ranging from 40 to 200 nm and a small fraction of large particles with an average particle size of 385 nm ranging from 250 to 400 nm [28]. It is worth pointing out that the small particles at GBs are dominantly ZrC, while the large particles are W-Zr-C-O complexes, which eliminate the free O at GBs and purifying GB interfaces. Therefore, the third layer interface is the nanoscale or sub-micrometer scale interface. The intuitive magnifying HRTEM of PBs shown in **Figure 8a**–**e** exhibits a perfect coherent structure interface between W matrix and ZrC dispersoids. This atomic ordered interface could pin and accumulate dislocations and thus effectively raise the strength and simultaneously improve the ductility of alloys.

The multi-scaled interface structure is in favor of irradiation resistance and decreases the retention of hydrogen. Mother grain boundaries may provide rapid diffusion paths for H and its isotope [63], and thus H could easily diffuse to the surface of the specimen even at a relatively low temperature. Thus a large fraction of D atoms may escape from the W-0.5ZrC during the D plasma irradiation at 400 K, leading to low retention of D in materials. At the same time, the fine sub-grains increase the GB interface density, and the nanosized particles create a high density of PB interfaces, which can absorb interstitial defects and then annihilate nearby vacancies by re-emitting the interstitial atoms back into the grain, thus improving the ability of irradiation resistance of W-0.5ZrC alloy.

#### **Figure 8.**

*(a) HRTEM image of W matrix and ZrC phase (intragranular) as viewed along [001]. (b) The SAEDP revealing the particle with a face-centered cubic structure. (c) Fast Fourier transform (FFT) pattern of selected red square area A on ZrC. (d) FFT pattern of selected red square area B at interface area between W and ZrC. (e) It is clear that the particle-matrix phase boundaries have a coherent structure like that shown in high magnification [28].*
