Abstract:
Vanadium alloys and their fabrication to produce materials for fusion applications having small additions of Ti, C and Zr that improve resistance to helium embrittlement.

Description:
This invention relates to vanadium alloys for fusion applications and more particularly, to vanadium alloys having small amounts of Ti, C, and Zr added for improved performance in fusion applications and was developed pursuant to a contract with the U.S. Department of Energy. 
     BACKGROUND OF THE INVENTION 
     The selection of a suitable material for use as the first wall of a fusion reactor has been a continuing problem. One of the most serious problems associated with the first wall is radiation damage caused by the energetic neutrons emitted by the plasma. The neutrons cause irradiation hardening and also produce helium gas atoms within the matrix of a first wall material which collects on the grain boundaries of most structural alloys. 
     For this application, vanadium alloys have certain advantages compared with conventional structural materials such as stainless steels because they exhibit low residual activity after irradiation, high thermal conductivity and low thermal expansion. They also have good mechanical strength at high temperatures and good corrosion resistance when used with lithium, which has been proposed as a first wall coolant. However, these alloys are subject to radiation induced degradation of mechanical properties that comes from matrix hardening and from grain boundary embrittlement due to helium accumulation at the grain boundaries. Therefore, there is a continuing need to improve the resistance of vanadium alloys to these radiation-induced phenomena. 
     SUMMARY OF THE INVENTION 
     In view of the above needs, it is an object of this invention to provide an alloy for fusion applications that resists helium embrittlement. 
     It is another object of this invention to provide an alloy for fusion applications that resists radiation hardening. 
     An additional object of this invention is to provide a vanadium alloy which when subjected to suitable thermomechanical treatment processes, forms a dispersion of MC type carbide particles that improves the performance of the alloy for fusion applications. 
     Additional objects, advantages and novel features of the invention will be set forth in part in the description which follows, and in part will become apparent to those skilled in the art upon examination of the following or may be learned by practice of the invention. The objects and advantages of the invention may be realized and attained by means of the instrumentalities and combinations particularly pointed out in the appended claims. 
     To achieve the foregoing and other objects and in accordance with the purpose of the present invention, as embodied and broadly described herein, the composition of this invention may consist essentially of vanadium and a sufficient amount of Ti and C to cause formation of a large number of small MC-type particles uniformly throughout the material. In the preferred embodiment the alloy comprises from 93.5 to 99.45 wt % V, from 0.5 to 6 wt % Ti and from 0.05 to 0.5 wt % C. In addition the invention consists essentially of vanadium and a sufficient amount of Ti, C and Zr to cause formation of two MC-type precipitates with different compositions, each to be distributed uniformly throughout the material. In the preferred embodiment the alloy comprises from 93 to 99.35 wt % V, from 0.5 to 6 wt % Ti, from 0.05 to 0.5 wt % C and from 0.1 to 0.5 wt % Zr. When these alloys are subjected to proper thermomechanical treatment, small, thin disks of MC are formed, providing improved properties that make them useful in fusion environments. 
    
    
     DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENT 
     The invention is comprised of vanadium alloys having dispersions of tiny face-centered-cubic (fcc) carbide particles. Alloys having only V, Ti and C contains a single MC-type carbide particle, whereas the alloys having Zr added contain two MC-type carbides with two separate and distinctly different compositions. The compositions of the alloys are in the Table. 
     
                       TABLE______________________________________      Composition, wt %Alloy No.    V      Ti        Zr   C______________________________________V1           98.45  1.5            0.05V2           98.40  1.5            0.10V3           98.25  1.5            0.25V4           96.90  3.0            0.10V7           98.30  1.5       0.15 0.05V8           98.25  1.5       0.15 0.10 V10         96.75  3.0       0.15 0.10 V11         98.10  1.5       0.30 0.10______________________________________ 
    
     Ingots weighting 400 g were arc-cast under argon into water-cooled copper molds. The cast microstructure for each alloy had elongated particles of carbide. The size, from 1 to 6 μm long, and number of the particles generally increase with the carbon content. Analytical transmission electron microscopy (TEM) revealed that many more, even smaller carbides, less than 0.2 μm, exist in the as cast microstructure. The V-Ti-C alloys contain small, thin disks of MC where M is about 25 wt % V and 75 wt % Ti. The V-Ti-Zr-C alloys also contain MC but M is composed of approximately 13.5 wt % V, 70.9 wt % Ti and 15.6 wt % Zr. A second, small (200 nm) MC-type carbide was observed in these alloys which tended to be blocky or geometric shaped. These particles are Zr rich with M equal to about 1.6 wt % V, 18.3 wt % Ti and 80.1 wt % Zr. 
     The composition and process described in the following example is intended to be illustrative and not in any way a limitation on the scope of the invention. Persons of ordinary skill in the art should be able to envision variations on the general principle of this invention that fall within the scope of the generic claims the follow. 
     EXAMPLE 
     The cast ingots were solution annealed at 1200° C. for 1 hour in a vacuum furnace (p&lt;10 -4 )Pa) to dissolve most of the MC particles. The ingots were warm rolled at 500° C. from a starting thickness of 1.27 cm down to 0.76 mm with intermediate 30 minute anneals at 1000° C. after reaching a thickness of 0.57 and 0.13 cm. Prior to the intermediate anneals and after the final rolling passes, the vanadium alloy sheets were pickled in solution consisting of 6 parts H 2  O, 3 parts HNO 3 , and 1 part HF by volume. The sheets were wrapped in Cb-1Zr foil to serve as a getter for the intermediate anneals which were conducted in a vacuum furnace (p&lt;10 -3  Pa). The rolling of the alloys having lower carbon content was done with relative ease as evidenced by the smooth surfaces and edges of the final fabricated sheets. However as more carbon is added, the alloy became stronger and more difficult to roll. When the alloy is strengthened too much the sheets can crack during the rolling process unless the fabrication parameters are properly adjusted. 
     Miniature tensile specimens were machined from the 0.76 mm sheet. Disks 3 mm in diameter for TEM were punched from the sheet and ground down to about 0.3 mm. Tensile specimens and TEM disks from all of the alloys were solution annealed at 1250° C. for 1 hour. These are referred to as &#34;SA&#34; specimens. Some of the SA specimens of both types were subsequently aged in quartz capsules under 1/2 atm of argon at 800° C. for two weeks; this condition was called &#34;SA+A&#34;, or solution annealed and aged. Finally a third group of as-rolled or 50% warm-worked specimens were also thermally aged using similar parameters to create the &#34;WW+A&#34;, or warm-worked at 500° C. and aged condition. The microstructure for the V-Ti-C and V-Ti-Zr-C, alloys VI and V11, produced by the heat treatments were similar. That is, the SA condition leaves the microstructure with only a few, but sometimes large precipitates. Aging at 800° C. (SA+A) precipitates thin disks of the Ti-rich MC-type carbides. The disks are parallel to the (100) planes in the crystal and, therefore, have three variants or orientations within each grain. The WW+A treatment creates a cell structure with many subgrain boundaries and dislocation segments. The carbide particles produced by this latter treatment range from 100 to 200 nm in size and have an equiaxed morphology instead of the oriented thin disks found in the SA+A condition. The heat treatment is important since the carbide-toughened alloys can be easily fabricated in the &#34;soft&#34; condition and subsequently heat treated to precipitate out the carbides and generate the desired tensile properties. 
     The tensile properties for the eight advanced alloys are listed in Table 2. 
     
                                           TABLE 2__________________________________________________________________________Tensile properties          TestSpecimen       Temperature                 Stress. MPa                          Elongation, %Alloy    No.  Condition          (°C.)                 Yield                     Ultimate                          Uniform                               Total__________________________________________________________________________V1  V107 SA.sup.a          420    266 438  11.2 17.0V1  V108 SA    520    255 515  18.1 22.8V1  V104 SA    600    256 524  16.6 22.3V1  V111 SA + A.sup.b          420    313 424  8.0  13.3V1  V113 SA + A          520    327 438  7.3  14.7V1  V121 SA + A          600    312 429  8.7  16.7V1  V103 WW + A.sup.c          420    353 412  5.7  13.3V1  V106 WW + A          520    356 424  6.3  14.7V1  V118 WW + A          600    366 429  5.3  13.0V2  V202 SA    420    242 521  14.8 20.6V2  V215 SA    520    240 442  10.8 15.6V2  V217 SA    600    251 582  18.3 22.6V2  V219 SA + A          420    325 475  8.3  15.0V2  V223 SA + A          520    346 492  8.7  16.0V2  V224 SA + A          600    331 463  8.7  16.7V2  V204 WW + A          420    369 415  6.0  14.0V2  V212 WW + A          520    369 440  6.7  16.7V2  V221 WW + A          600    383 457  6.7  14.3V3  V309 SA    420    237 424  10.8 16.2V3  V317 SA    520    237 538  14.8 21.0V3  V313 SA    600    237 570  20.0 30.5V3  V307 SA + A          420    293 482  9.0  14.3V3  V304 SA + A          520    336 589  11.7 16.3V3  V315 SA + A          600    308 578  12.3 25.0V3  V302 WW + A          420    461 586  9.7  13.3V3  V305 WW + A          520    475 593  9.7  13.0V3  V321 WW + A          600    458 593  8.7  21.7V4  V410 SA    420    324 404  13.8 22.3V4  V418 SA    520    284 551  19.6 27.0V4  V420 SA    600    281 571  20.3 28.0V4  V402 SA + A          420    219 376  11.0 18.7V4  V411 SA + A          520    221 391  11.6 19.5V4  V421 SA + A          600    224 388  14.2 19.0V4  V405 WW + A          420    293 376  8.3  15.9V4  V407 WW + A          520    280 372  9.1  16.8V4  V412 WW + A          600    296 393  7.5  13.8V7  V713 SA    420    246 414  12.2 20.5V7  V714 SA    520    240 520  15.9 22.0V7  V721 SA    600    244 567  15.8 21.6V7  V703 SA + A          420    250 394  10.3 19.0V7  V704 SA + A          520    263 416  10.3 18.0V7  V706 SA + A          600    253 406  11.0 16.7V7  V702 WW + A          420    335 394  6.7  15.0V7  V705 WW + A          520    339 409  7.0  14.0V7  V722 WW + A          600    342 407  7.3  12.7V8  V801 SA    420    226 494  13.8 18.8V8  V804 SA    520    207 518  15.5 20.9V8  V817 SA    600    228 569  20.6 27.6V8  V802 SA + A          420    220 402  10.0 16.0V8  V809 SA + A          520    240 420  10.3 17.3V8  V818 SA + A          600    234 431  12.7 21.7V8  V811 WW + A          420    349 419  7.0  14.3V8  V820 WW + A          520    362 432  6.7  12.3V8  V821 WW + A          600    374 446  7.0  15.0 V10V1001    SA    420    245 432  13.2 22.4 V10V1019    SA    520    246 533  15.5 23.0 V10V1015    SA    600    237 511  16.2 25.4 V10V1003    SA + A          420    199 344  11.7 21.0 V10V1007    SA + A          520    188 350  12.3 20.0 V10V1008    SA + A          600    201 357  11.3 19.7 V10V1010    WW + A          420    296 381  8.0  15.7 V10V1011    WW + A          520    293 372  8.7  16.0 V10V1013    WW + A          600    305 376  7.3  13.7 V11V1123    SA    420    244 433  14.0 22.5 V11V1105    SA    520    254 558  15.0 20.6 V11V1119    SA    600    247 568  19.4 26.2 V11V1110    SA + A          420    317 456  8.9  17.0 V11V1122    SA + A          520    324 466  8.0  16.0 V11V1124    SA + A          600    319 472  9.3  17.7 V11V1104    WW + A          420    388 456  5.7  11.3 V11V1108    WW + A          520    408 483  5.7  11.7 V11V1113    WW + A          600    403 469  5.7  12.0__________________________________________________________________________ .sup.a SA = solution annealed 1 h at 1250° C. .sup.b SA + A = solution annealed 1 h at 1250° C. and aged two weeks at 800° C. .sup.c WW + A = 50% warm worked and aged two weeks at 800° C. 
    
     Looking at the results for VI, the SA+A treatment produces a stronger alloy, having higher yield strength, with less ductility, total elongation, than the SA treatment. The WW+A has even higher yield strengths with slightly lower ductility compared to the SA+A. These results would be expected from the microstructures of the respective alloys. The results are generally the same for the other alloys with minor differences depending on the amount of carbon and whether Zr is added or not. 
     Additional carbon, above the lowest level of 0.05 wt %, did not always make the alloys stronger because the carbide particles were coarsened and their numbers reduced. Thus, precipitation hardening effects were minimized. In some cases, the addition of Zr provided additional increases in ultimate strength with little or no loss in ductility as shown by V7-SA compared with V1-SA. In other cases the alloys with the Zr additions were slightly weaker and more ductile as seen by comparing V8 and V2. 
     Increasing the amount of Ti from 1.5 to 3.0 wt % increased the strength and ductility of V2 to that of V4, but the gain in strength was lost upon aging. This is because the solid solution strengthening offered by the Ti in the SA condition was lost when the Ti precipitated out in the Ti-rich MC particles. What resulted was a weaker matrix with MC particles that were too coarse to add substantial precipitation-hardening. 
     The specimens were also tested for resistance to helium embrittlement. All four V-Ti-C alloys increased their ductility at 420° C. with addition of helium and exhibited only modest losses as test temperatures increased. The V-Ti-Zr-C alloys showed similar behavior with the exception of V7 and V10 which had slight losses at 420° C. rather than increases. This represents a dramatic improvement in resistance to helium embrittlement relative to other vanadium alloys. The reason for the improved resistance to helium embrittlement demonstrated by the advanced V-Ti-C and V-Ti-Zr-C alloys is that thin Ti-rich MC disks in the alloys act as sinks for the helium in the microstructure. This prevents the helium from migrating to the grain boundaries, which causes embrittlement. In addition, Ti-rich MC particles in the grain boundaries themselves directly trap helium that flows to the grain boundaries. These trapping mechanisms enable the alloys to accommodate more helium before serious helium embrittlement occurs. The WW+A condition also is effective in resisting helium embrittlement, but for different reasons. With the warm-worked microstructure, helium is trapped on dislocations, at subgrain boundaries, and at equiaxed MC particles. Furthermore there appear to be no high angle grain boundaries in the WW+A material where helium collects to form the largest bubbles in the microstructure.