**2. γ‐UMo massive hydride discovery and powder production**

Several methods have been used to comminute UMo alloys: the more sophisticated centrifugal atomization already mentioned from KAERI, mechanical grinding performed initially in Canada and rotating electrode [36]. Studies of gaseous atomization with gold used as surrogate material to evaluate production performance were also performed [37]. Another explored possibility of comminuting UMo ductile alloys is by the decomposition of metastable γ‐UMo in α‐U and U2Mo by a controlled heat treatment. When an incipient cellular precipitation totally decorates grain boundaries, the α‐U phase present can be hydrided. The brittle hydride weakens the material that can easily be milled to particles of dimensions similar to the original grain size [38, 39].

Increasing the temperature of a quenched melted rod of U7Mo in a hydrogen atmosphere, hydrogen pick up started at around 100°C; the incorporation went higher than the hydrogen solubility limit and before decomposition of the metastable gamma phase took place [40,41]. It could be observed that a new phase begun to form at the place where tensile stresses were bigger. With this input, the method to completely hydride the γ‐UMo alloy was developed. An hydriding, milling and dehydriding method, HMD, was set forward to obtain γ‐UMo powder [42, 43].

## **2.1. Casting of UMo with stress retention**

case of uranium alloy worked as a monolithic meat with UAl4 precipitates, the uranium loading must be less than 35 wt.% uranium (4 at.% U)—to avoid inhomogeneous dispersions that could produce hot spots—rendering a final meat density of 1.35 gU/cm3 of total uranium. In the case of a UAlx (x = 2, 3 or 4) dispersed powder in an aluminum matrix with a concentration between 40 and 50 v/v%, densities of total uranium in the meat can reach values higher than 2 gU/cm3

One of the simplest choices to increase fuel density was to develop fuel plates with LEU U3O8 dispersed powder in an aluminum matrix. The density of U3O8 is 8.3 g/cm3 and the meat total

aluminum matrix is less stable than U3O8 since important swelling is formed due to low density

Uranium silicide compounds such as U3Si and U3Si2 have different behavior under irradiation. While the first one presented in some cases break away swelling [18], U3Si2 was finally selected to rich total uranium densities at the fuel plate meat of 4.8 gU/cm3 [19]. This fuel is the last one that has been qualified for general uses in research reactors using LEU. It is worthwhile to comment that U3Si is a ductile compound and centrifugal atomization [20] was specially

By 1996, in a systematic study, several high density uranium compounds were revisited to be used as nuclear fuels with the intention of reducing even more the utilization of HEU in research reactor fuels and to look for a fuel more easily reprocessed than the U3Si2 dispersions [21, 22]. Some of these compounds to be qualified resulted to be gamma stabilized uranium alloys [23], in particular the metastable phase γ‐UxMo. Quickly, it was determined in irradi‐ ation experiments that the needed composition had to have more than 6% w/w Mo (x > 6) [24]; afterwards it was observed that γ‐UMo had an exceptional behavior in the allocation of fission gas bubbles [25]. UxMo alloys, with weights percent between seven and ten (7 ≤ x ≤ 10), are being tested since the stability of the gamma phase is enhanced as molybdenum concentration increases [26, 27], favoring intermediate fabrication processes. UMo presents interface incompatibilities with aluminum at high irradiation fluxes when fission gases generate undesired porosity [28–30]; this drawback is diminished by the incorporation of silicon to the aluminum matrix [31]. Fuel/cladding interaction can be minimized by covering the fuel particles with a diffusion barrier material or tailoring the fuel or matrix materials with the

Since γ‐UMo is a ductile alloy, it can be used as a monolithic fuel; nevertheless it cannot be colaminated with aluminum because of the different thermomechanical properties of both materials [32]. Several alternatives have been proposed to obtain a γ‐UMo monolithic fuel with aluminum cladding [33]. They comprise a first step in which UMo foils are hot laminated to final dimensions and a second step is the incorporation of the aluminum cladding that can be performed by transient liquid phase bonding, friction stir welding or hot isostatic pressing. All of these three alternatives have UMo/Al interfaces where fission gas bubbles can coales‐ cence. To avoid a UMo/Al interface a zirconium diffusion barrier can be incorporated between

uranium density can reach values higher than 3.1 gU/cm3

reaction products and diffusion porosity formed [17].

92 Nuclear Material Performance

dedicated to obtain a ductile compound in powder form.

**1.3. Uranium molybdenum long run qualification**

incorporation of additional alloying elements.

.

. Dispersed UO2 powder in an

Batches of up to one kilogram of alloy were prepared with natural uranium (>99.7% U) and LEU (>99.9% U). Lumps of metallic uranium and molybdenum chips were used to prepare a 7% w/w molybdenum alloy. Uranium was deoxidized using nitric acid and molybdenum with melted sodium hydroxide. A high frequency induction furnace was used for melting the alloy in magnesium oxide crucibles and an inert gas atmosphere.

The melt was poured in a graphite dish obtaining a circular plate of approximately 8 mm thick and 120 mm diameter. Tension stresses at the top of the plate are retained since solidification begins at the bottom in contact with the graphite. The plate is finally broken in pieces (**Figure 1**) such that they can be incorporated in a hydriding chamber.

Residual tension stresses at the upper surface were checked by peak displacement in X‐ray diffractions (DRX) in the as melted condition and after an annealing treatment (**Figure 2**). These stresses in the γ‐UMo alloy are needed to allow the initiation of hydriding.

**Figure 1.** Eight millimeters thick uranium molybdenum lumps with residual stresses.

**Figure 2.** Evidence of stresses in γ‐U7Mo shown by the displacement of the (011) DRX peak of a just melted and an annealed sample.

#### **2.2. Low temperature hydriding of γ‐UMo alloy**

γ‐U7Mo alloys were hydrided in a 1.5 l chamber vacuum and temperature assisted. A flow meter was used to measure the incoming hydrogen needed to keep a constant hydrogen pressure during absorption. In the stressed sample, hydrogen can be allocated in interstitial positions in the crystal structure (solubilization) or in hydrogen traps in dislocations generated by tension stresses. Traps must be filled with hydrogen to allow massive hydriding of γ‐UMo afterwards. At 230°C, in a 10-3 torr vacuum, high purity hydrogen is introduced up to a pressure of 1 atmosphere, filling traps of the stressed γ‐U7Mo alloy with some hundreds ppm of hydrogen in less than one hour [44]. Lowering the temperature to 120°C, hydrogen absorption begins after a few minutes; typical hydrogen absorption rates can reach values higher than 1 liter of hydrogen gas—at standard temperature and pressure—per minute per kilogram of alloy. After several hours, hydriding is completed. Hydrogen absorption dependence with time is shown in **Figure 3**. Additional temperature controls are needed during the process since the hydriding reaction is exothermic.

**Figure 3.** γ‐U7Mo time dependence absorption of gaseous hydrogen at 120°C.

**Figure 1.** Eight millimeters thick uranium molybdenum lumps with residual stresses.

annealed sample.

94 Nuclear Material Performance

**2.2. Low temperature hydriding of γ‐UMo alloy**

the hydriding reaction is exothermic.

**Figure 2.** Evidence of stresses in γ‐U7Mo shown by the displacement of the (011) DRX peak of a just melted and an

γ‐U7Mo alloys were hydrided in a 1.5 l chamber vacuum and temperature assisted. A flow meter was used to measure the incoming hydrogen needed to keep a constant hydrogen pressure during absorption. In the stressed sample, hydrogen can be allocated in interstitial positions in the crystal structure (solubilization) or in hydrogen traps in dislocations generated by tension stresses. Traps must be filled with hydrogen to allow massive hydriding of γ‐UMo afterwards. At 230°C, in a 10-3 torr vacuum, high purity hydrogen is introduced up to a pressure of 1 atmosphere, filling traps of the stressed γ‐U7Mo alloy with some hundreds ppm of hydrogen in less than one hour [44]. Lowering the temperature to 120°C, hydrogen absorption begins after a few minutes; typical hydrogen absorption rates can reach values higher than 1 liter of hydrogen gas—at standard temperature and pressure—per minute per kilogram of alloy. After several hours, hydriding is completed. Hydrogen absorption dependence with time is shown in **Figure 3**. Additional temperature controls are needed during the process since

Optimum hydriding conditions were set up by observing the different hydrogen pick up rates in an increasing temperature slope and in a decreasing one, after hydrogen saturation of traps at 230°C. In **Figure 4**, it can be observed that hydrogen absorption rate while heating has a maximum value of 220 ppm/l and while cooling is much higher, 650 ppm/l, for a 700 g batch. From this curves it was determined that hydriding conditions were between 50 and 190°C, and the maximum rate is at 120°C.

**Figure 4.** Hydrogen absorption rates during heating and cooling of a γ‐U7Mo batch.

If a piece of γ‐UMo is heated above 230°C in a hydrogen atmosphere, tension stresses will appear as it is cooled, inducing the massive hydriding of the material. This explains why in a cooling cycle hydrogen absorption is higher than in a heating one [42]. U10Mo has also been hydrided with similar results utilizing selected heating and cooling cycles such that hydriding conditions can be reached without initially incorporating residual stresses [45, 46].

After the hydride is totally formed at 120°C, it is convenient to increase the temperature up to 325°C and evacuate the chamber to eliminate the hydrogen in traps; this procedure makes the hydride less pyrophoric. Finally, at room temperature, air is introduced in a temperature controlled way such that the hydride is passivated.

#### **2.3. Characterization of the γ‐UMo hydride**

The γ‐U7Mo hydride is dark gray, brittle, fragmented in platelets, with small transgranular cracks (**Figure 5**). It is pyrophoric and burns with flame because of hydrogen liberation. When oxidized in air, it is dark brown. Rietveld refinement of a X‐Ray diffraction—XRD—pattern (**Figure 6**) show wide peaks corresponding to a unique stressed crystalline cubic *A‐15* structure (space group *Pm*3 ¯*n*, *N° 223*) of the β‐W prototype, the same as β‐UH3 [47, 48] with eight heavy atoms in its unit cubic cell with parameter 6.6598 Å; the stoichiometry is of (U7Mo)H3‐y with values of *y* smaller than 0.2 obtained experimentally by weight difference. The XRD hydride density is 10.39 g/cm3 . Since the density of γ‐U7Mo with a unit bcc cell parameter of 3.4785 Å is 17.5 g/cm3 , the increase in volume during hydriding is 68%, causing the fragmentation of the brittle hydride that is being formed.

**Figure 5.** γ‐U7Mo hydrided fragments.

**Figure 6.** XRD, Rietveld refinement and indexation of (U7Mo)H3 A‐15 structure.

#### **2.4. Milling of the γ‐UMo hydride**

Vickers hardness of the hydride (U7Mo)H3 is approximately 300 VH. This hydride must be milled in a sufficient inert atmosphere to avoid burning up. Low impact mills are preferred to avoid excess fines (particle size smaller than 45 μm). Two roll mills and/or conical vibratory crushers were used (**Figure 7**). Hydride was first reduced with a manual roll mill in a glove box with less than 5% oxygen (mesh #10, 2 mm opening). These particles were then passed once through a conical crusher in a dynamic inert atmosphere reducing the size (#120, 125 microns). **Figure 8** shows hydride (γ‐U7Mo)H3 milled particles of 80 μm mean size. Passivation is needed before exposing the particles to air.

Gamma Uranium Molybdenum Alloy: Its Hydride and Performance http://dx.doi.org/10.5772/63652 97

**Figure 7.** Low impact mills: roll (left) and gyratory cone (right).

(**Figure 6**) show wide peaks corresponding to a unique stressed crystalline cubic *A‐15* structure

atoms in its unit cubic cell with parameter 6.6598 Å; the stoichiometry is of (U7Mo)H3‐y with values of *y* smaller than 0.2 obtained experimentally by weight difference. The XRD hydride

¯*n*, *N° 223*) of the β‐W prototype, the same as β‐UH3 [47, 48] with eight heavy

, the increase in volume during hydriding is 68%, causing the fragmentation of

. Since the density of γ‐U7Mo with a unit bcc cell parameter of 3.4785 Å

(space group *Pm*3

96 Nuclear Material Performance

is 17.5 g/cm3

density is 10.39 g/cm3

the brittle hydride that is being formed.

**Figure 5.** γ‐U7Mo hydrided fragments.

**2.4. Milling of the γ‐UMo hydride**

is needed before exposing the particles to air.

**Figure 6.** XRD, Rietveld refinement and indexation of (U7Mo)H3 A‐15 structure.

Vickers hardness of the hydride (U7Mo)H3 is approximately 300 VH. This hydride must be milled in a sufficient inert atmosphere to avoid burning up. Low impact mills are preferred to avoid excess fines (particle size smaller than 45 μm). Two roll mills and/or conical vibratory crushers were used (**Figure 7**). Hydride was first reduced with a manual roll mill in a glove box with less than 5% oxygen (mesh #10, 2 mm opening). These particles were then passed once through a conical crusher in a dynamic inert atmosphere reducing the size (#120, 125 microns). **Figure 8** shows hydride (γ‐U7Mo)H3 milled particles of 80 μm mean size. Passivation

**Figure 8.** (U‐7Mo)H3 powder of 80 μm mean size. SEM.

## **2.5. (U7Mo)H3 dehydriding**

The dehydriding of the powder is done in a vacuum atmosphere with a decrease of hydride size that is important because of the density difference with the final γ‐U7Mo powder. Any gas evolution is followed by pressure measurements with a closed chamber; periodic evacu‐ ation is performed to reduce excess pressure. Hydrogen liberation begins at 125°C and beyond 425°C can be bursting. The hydrogen in traps—that is incorporated at 230°C—is liberated at temperatures between 300 and 375°C, whereas the hydrogen in interstitial sites—incorporated at 120°C—will begin to be evacuated at temperatures above 380°C as shown in **Figure 9**. If, after hydriding, hydrogen in traps is removed as mentioned in 2.2., the bump of hydrogen pressure increase at 325°C will not be observed.

**Figure 9.** Hydrogen evolution during dehydriding. Hydrogen in traps and interstices are liberated around 325°C and beyond 425°C, respectively.

To eliminate hydrogen to lower values than 50 ppm heating up to 700°C in a dynamic vacuum atmosphere during 2 hours is needed. The heating chamber was vibrated to avoid the sintering of the particles. When the dehydriding is finished, an inert gas is introduced and quick cooling to room temperature is needed to quench the metastable γ‐U7Mo phase. Controlled passiva‐ tion is done by slow removal of the inert gas and air introduction. Dehydriding needs at least 5 hours. Cracks formed during the hydriding process are still present after dehydriding (**Figure 10**). The powder must be kept in an inert atmosphere to avoid oxidation.

**Figure 10.** HMD γ‐U7Mo powder. Big particles are 100 μm size (SEM).

If it is of interest, cracks can be sintered with a heat treatment of 8 hours at 1000°C in a vacuum atmosphere in the vibrating chamber [44]. **Figure 11** shows the results of this treatment.

**Figure 11.** SEM (first row) and metallographies (second row) of γ‐U7Mo powder before (left column) and after (right column) heat treatment at 1000 °C in a vacuum atmosphere. Big particles are 100 μm.

#### **2.6. Absorption, coverage and lamination**

after hydriding, hydrogen in traps is removed as mentioned in 2.2., the bump of hydrogen

**Figure 9.** Hydrogen evolution during dehydriding. Hydrogen in traps and interstices are liberated around 325°C and

To eliminate hydrogen to lower values than 50 ppm heating up to 700°C in a dynamic vacuum atmosphere during 2 hours is needed. The heating chamber was vibrated to avoid the sintering of the particles. When the dehydriding is finished, an inert gas is introduced and quick cooling to room temperature is needed to quench the metastable γ‐U7Mo phase. Controlled passiva‐ tion is done by slow removal of the inert gas and air introduction. Dehydriding needs at least 5 hours. Cracks formed during the hydriding process are still present after dehydriding

If it is of interest, cracks can be sintered with a heat treatment of 8 hours at 1000°C in a vacuum atmosphere in the vibrating chamber [44]. **Figure 11** shows the results of this treatment.

(**Figure 10**). The powder must be kept in an inert atmosphere to avoid oxidation.

**Figure 10.** HMD γ‐U7Mo powder. Big particles are 100 μm size (SEM).

pressure increase at 325°C will not be observed.

beyond 425°C, respectively.

98 Nuclear Material Performance

The surface absorption of HMD powder in an air atmosphere was studied by increasing the temperature in a closed vibrating chamber. The volume of the chamber and piping was of approximately 3 liters and runs with 64 and 130 grams of γ‐U7Mo HMD powder were performed [49]. **Figure 12** shows the pressure evolution from an initial value of 800 mbar air atmosphere as temperature is increased up to 500°C. The initial increment in pressure corresponds to surface gas desorption, fundamentally water at 120°C. A first run with 64 g showed an important pressure reduction up to a value of 200 mbar. Two consecutive runs were performed with a fresh 130 g sample of γ‐U7Mo to evaluate if equilibrium is reached or a diffusion barrier was formed. The first run with this second batch, curve 130 g in **Figure 12**, was heated up to equilibrium. The second run, curve 130 g bis, was performed after cooling and reintroducing air in the chamber. These last two runs showed that pressure equilibrium was achieved without reaching saturation. Oxygen and nitrogen were incorporated without presenting a barrier to gas diffusion; nitrogen incorporation is known that begins to be important at temperatures higher than 300°C in metallic uranium [40]. A similar scenario will be present inside a picture and frame ensemble before hot rolling.

**Figure 12.** Pressure evolution with temperature of γ‐U7Mo HMD powder exposed to an air atmosphere in a closed chamber.

Coverage of γ‐UMo powder is of interest not only to diminish oxidation during storage but fundamentally to have a diffusion or metallurgical barrier protection against undesirable growth of, for example, the UMo/Al interaction zone during high flux irradiation. Surrogate powders were used to initially set up equipment and coverage conditions; silicon, aluminum, magnesium and silicon were used with chemical vapor deposition and dip coating techniques [50, 51]. Physical vapor deposition [52] is another possibility of coverage technique. Mixed powders of UMo and silicon in the vibrating chamber at 950°C in vacuum during 2 hours cover the bigger UMo particles with silicon [51].

The coverage of ductile UMo particles with brittle materials is thermomechanically incom‐ patible in layers thicker than 1 micron. Probably ductile nickel and niobium seem to be better candidates as they have been tested previously in other nuclear fuel developments. New technologies are now available that consist in particles spheroidization and coverage using an induction couple plasma equipment that nowadays can be purchased in a commercial form [53].

The brittle transgranular fractures produced during comminution of the hydride gives as result polyhedral shaped particles. **Figure 13** shows a metallography of the laminated meat of a dispersed HMD miniplate with γ‐U7Mo powder in an aluminum matrix. The irregular shape of HMD particles favors the mixing with Al powder, diminishes flow segregation—compared with spherical particles—and green compacts practically have no shredding. Cracks in the HMD particles do not propagate during hot lamination since UMo is a ductile alloy. Final porosity of a meat with HMD powder is between 5 and 10% v/v. Dispersed UMo particles in an aluminum matrix miniplates were also colaminated using silicon covered U7Mo HMD powder and centrifugal atomized particles with LEU and natural uranium [51].

**Figure 13.** Metallography of the meat of γ‐U7Mo HMD powder dispersed in an aluminum matrix of a laminated mini‐ plate. The size of the big particle is approximately 100 microns. Holes between particles were fundamentally produced during polishing.

Monolithic UMo meats can be obtained by powder metallurgy. A coupon of U7Mo HMD powder was cold pressed and surrounded by AISI 304L stainless steel in the standard lids and frame configuration to conform the cladding (**Figure 14**). Colamination was performed at 675°C, a temperature in the gamma phase stable zone of the U7Mo, in a nitrogen atmosphere. The use of a UMo monolithic meat elaborated by powder metallurgy allows an easier way of obtaining different meat widths in the same plate using conformational dies, incorporation of powdered neutron moderators such as high temperature stable hydrides, burnable poisons and also nanosized porous powders to adsorb fission gases at grain boundaries so as to reduce overall swelling of fuel plates [54].

**Figure 14.** Monolythic UMo miniplate with stainless steel cladding.

**Figure 12.** Pressure evolution with temperature of γ‐U7Mo HMD powder exposed to an air atmosphere in a closed

Coverage of γ‐UMo powder is of interest not only to diminish oxidation during storage but fundamentally to have a diffusion or metallurgical barrier protection against undesirable growth of, for example, the UMo/Al interaction zone during high flux irradiation. Surrogate powders were used to initially set up equipment and coverage conditions; silicon, aluminum, magnesium and silicon were used with chemical vapor deposition and dip coating techniques [50, 51]. Physical vapor deposition [52] is another possibility of coverage technique. Mixed powders of UMo and silicon in the vibrating chamber at 950°C in vacuum during 2 hours cover

The coverage of ductile UMo particles with brittle materials is thermomechanically incom‐ patible in layers thicker than 1 micron. Probably ductile nickel and niobium seem to be better candidates as they have been tested previously in other nuclear fuel developments. New technologies are now available that consist in particles spheroidization and coverage using an induction couple plasma equipment that nowadays can be purchased in a commercial form

The brittle transgranular fractures produced during comminution of the hydride gives as result polyhedral shaped particles. **Figure 13** shows a metallography of the laminated meat of a dispersed HMD miniplate with γ‐U7Mo powder in an aluminum matrix. The irregular shape of HMD particles favors the mixing with Al powder, diminishes flow segregation—compared with spherical particles—and green compacts practically have no shredding. Cracks in the HMD particles do not propagate during hot lamination since UMo is a ductile alloy. Final porosity of a meat with HMD powder is between 5 and 10% v/v. Dispersed UMo particles in an aluminum matrix miniplates were also colaminated using silicon covered U7Mo HMD

powder and centrifugal atomized particles with LEU and natural uranium [51].

chamber.

100 Nuclear Material Performance

[53].

the bigger UMo particles with silicon [51].
