**4. Discussion on irradiation-induced damage effects on mechanical properties of 5***xxx* **and 6***xxx* **series Al alloys**

Although fracture toughness data on irradiated Al alloys is scarce, significant data on tensile properties is available in the literature. In this section, tensile data on irradiated Al alloys collected from literature is plotted as a function of thermal fluence to understand the changes in tensile properties with the evolution of irradiation-induced microstructural damage (or transmutation-produced Si content). Once this relation is established, then one can make a bridge to correlate these changes to corresponding changes in fracture toughness properties, where only limited data is published in the literature.

In case of control rod drive follower (CRDF) A-2 tubes of the High Flux Beam Reactor (HFBR) in Brookhaven National Laboratory, USA, produced from 6061-T6 alloy, irradiated at 338 K

amorphous Si-rich particles are observed in the microstructure in place of original Mg2Si

compared to the HFR hotspot TFR value of maximum 1.4. The total measured Si at this fluence

The location of this transmutation-produced Si precipitates in the microstructure will have substantial impact on the mechanical properties of the alloys. In 1100 and 6061 alloys, it was identified that the transmutation-produced Si will precipitate as elemental Si particles, which are uniformly distributed in the matrix and associated with voids [9]. Farrell et al. [17] reported a noncrystalline Si-coating inside the voids of 1100-O Al alloy at a high thermal fluence (*E* <

decoration of original Mg2Si precipitates with transmutation-produced Si in addition to the association of Si particles with voids [9]. Precipitation of this Si along the grain boundary can lower the fracture toughness. For example, CRDF A-2 tubes of HFBR produced from 6061-T6 alloy have shown a drop in fracture toughness to ~8 (MPa)·m1/2 from an unirradiated value of

**Figure 4**). The microstructure of this alloy, with a very high transmutation-produced Si content of 8 wt.%, has shown large silicon flakes occupying less than one-fifth of the grain boundary area [5]. Similarly, heavy discontinuous precipitation at grain boundaries is observed in 5052

From the above discussion, it can be concluded that the transmutation-produced Si is the dominant irradiation damage mechanism in 5*xxx* and 6*xxx* series Al alloys irradiated at temperatures <373 K. Consequently, transmutation-produced Si is taken as the measure of the irradiation damage in HFR vessel wall. There are differences in how this transmutationproduced Si will influence the mechanical properties of 5*xxx* and 6*xxx* Al alloys, which will be

**4. Discussion on irradiation-induced damage effects on mechanical**

Although fracture toughness data on irradiated Al alloys is scarce, significant data on tensile properties is available in the literature. In this section, tensile data on irradiated Al alloys collected from literature is plotted as a function of thermal fluence to understand the changes in tensile properties with the evolution of irradiation-induced microstructural damage (or transmutation-produced Si content). Once this relation is established, then one can make a bridge to correlate these changes to corresponding changes in fracture toughness properties,

21.75 (MPa)·m1/2 after irradiation to a thermal neutron fluence of ~42 × 1026 n/m2

. The 6061 alloy irradiated to ~1027 n/m2 at ~328 K has shown a

[3].

, a high concentration of very fine (8 nm)

, which gives a high TFR of 21

at 338 K (see

up to a very high thermal fluence of 42 × 1026 n/m2

0.025 eV) of ~2.3 × 1027 n/m2

398 Radiation Effects in Materials

discussed in the next section.

precipitates [5]. The corresponding fast fluence is 2 × 1026 n/m2

was found to be ~8 wt.%, including 0.6% of the initial Si content.

alloy irradiated up to a thermal fluence of ~31 × 1026 n/m2

**properties of 5***xxx* **and 6***xxx* **series Al alloys**

where only limited data is published in the literature.

**Figure 1.** Literature tensile data of irradiated Al alloys in comparison with HFR SURP data. (a) Yield strength versus thermal fluence, (b) tensile strength versus thermal fluence and (c) total elongation versus thermal fluence.

**Figure 2.** Schematic diagram showing various regimes in irradiation hardening behavior of (a) 5xxx and (b) 6xxx series Al alloys.

Farrell et al. [3] published data on tensile behavior of 5052-O aluminum alloy (Al–2.2% Mg) heavily irradiated in HFIR to fluences greater than 1027 n/m2 in contact with cooling water at 328 K (see **Figure 1**). HFIR is predominantly a thermal reactor with a strong fast neutron component. The thermal neutron fluence (*E* < 0.0025 eV) of the samples ranged up to ~31 × 1026 n/m2 to produce 7.15 wt.% Si. The fast neutron fluence (*E* > 0.1 MeV) is a factor of 1.7 (=TFR) lower than the thermal fluence, and results in a damage value of 260 displacements per atom (dpa). Associated gas generation due to fast neutrons was estimated to be 8.5 × 10−5 atomic fraction He and 5 × 10−4 atomic fraction H. It is important to notice that this data is very relevant to the HFR SURP program because of the (i) similar chemical composition of 5052 and 5154 alloys, except that Mg in 5154-HFR alloy is 3.2% instead of 2.2% in 5052 and (ii) similar irradiation conditions, including temperature, TFR, and high fluence values. Farrell et al. [2] also published tensile data on heavily irradiated 6061-O and 6061-T6 alloys irradiated at 328 K up to ~31 × 1026 n/m2 and tested at three different temperatures (323, 373 and 423 K). **Figure 1** shows only results at 323 K due to their relevance to HFR operating conditions. Additionally, tensile strength data from a 6061-T6 type alloy tested from CRDF A-2 tubes of HFBR, published by Weeks et al. [5] are also shown in **Figure 1**.

Comparison of yield and tensile strength properties of all these alloys including HFR-SURP tensile data, shown with added trend lines in **Figure 1 (a, b)**, reveals specific trends in irradiation hardening and embrittlement behavior. Each of these alloys showed a rapid hardening regime (and corresponding drop in ductility) at the beginning, followed by a transition regime toward a relatively slow hardening (and stable ductility) regime. A brittle regime is observed in some alloys at the end, as shown schematically in **Figure 2**. Depending on whether an alloy is of 5*xxx* or 6*xxx* series, a single or multiple irradiation damage mechanism can be active, determining the hardening rates in each of these regimes.

#### **4.1. Tensile behavior of 6***xxx* **series alloys**

**Figure 2.** Schematic diagram showing various regimes in irradiation hardening behavior of (a) 5xxx and (b) 6xxx series

Farrell et al. [3] published data on tensile behavior of 5052-O aluminum alloy (Al–2.2% Mg) heavily irradiated in HFIR to fluences greater than 1027 n/m2 in contact with cooling water at 328 K (see **Figure 1**). HFIR is predominantly a thermal reactor with a strong fast neutron component. The thermal neutron fluence (*E* < 0.0025 eV) of the samples ranged up to ~31 ×

to produce 7.15 wt.% Si. The fast neutron fluence (*E* > 0.1 MeV) is a factor of 1.7 (=TFR)

Al alloys.

400 Radiation Effects in Materials

1026 n/m2

In case of 6*xxx* series Al alloys, rapid hardening observed at the onset of irradiation (regime 1), for example, curves of 6061-O and 6061-T6 alloys (**Figure 1 (a, b)**), can be attributed to irradiation-induced dislocation damage. It is known that the irradiation-induced dislocation density rapidly increases at the beginning and reaches a saturation value (see Section 3.1) at relatively low fluence values. For steels irradiated at 603 K (~0.35*Tm*), this saturation disloca‐ tion density is expected to reach around a fast fluence of ~2 × 1026 n/m2 (15). For Al alloys irra‐ diated at temperatures about 323 K (~0.35*Tm*), it is expected that saturation is reached at lower fluences due to the low effective displacement energy of Al (~25 eV) compared to Fe (~40 eV) [9].

In the transition regime (regime 2), precipitation of transmutation-produced Si takes over as the major contributing mechanism, while the dislocation density reaches a saturation limit. It is known from literature that the transmutation Si in 6*xxx* alloys nucleate as amorphous Si particles in the matrix, associate with irradiation-induced voids, and decorate existing Mg2Si precipitates (see Section 3.2). The contribution of this mechanism to irradiation hardening of 6*xxx* series alloys is low because deformation which occurs by the shearing of soft Si parti‐ cles produces little strain hardening [18]. At the same time, the hardening contribution from an increase in size of existing Mg2Si precipitates (due to Si decoration) is also low. This is because for a given particle density, the increase in precipitate size (*r*) and the decrease in planar spacing (*λ*) of the precipitates resulting from precipitate growth due to Si decoration have a minimal effect on strength based on the Orowan–Ashby equation [19] (Eq. (2)).

$$
\Delta \sigma = \frac{0.13Gb}{\lambda} \ln \frac{r}{b} \tag{2}
$$

Assuming a saturation density of ~ 6 × 1014 m−2 in Al alloys (same as in steel), a rough esti‐ mate of the total contribution of dislocation hardening can be made using the following equation [20]:

$$
\Delta \sigma\_{ds} = \sigma - \sigma\_o = Gb\rho^{1/2},\tag{3}
$$

where *σ* is the strength of the material after introducing dislocation structure, *σo* is the intrin‐ sic strength of the material with low dislocation density, *G* is shear modulus (=2.648 × 104 MPa), *b* is Burgers vector (0.286 nm) and *ρ* is dislocation density. The *G* and *b* values of pure Al taken from reference [21] are used here. Substituting *ρ* = 6 × 1014 m−2 into Eq. (3) gives Δ*σ* = ~186 MPa. The magnitude of hardening observed at the beginning of 6061-O and 6061-T6 alloys is in good agreement with this value. An estimation of the total irradiation hardening using the microstructure data (*r* = 4 nm, λ = 4 nm) of CRDF A-2 alloy at 42 × 1026 n/m2 ther‐ mal fluence taken from [5] resulted in an estimated yield strength of 605 MPa, which is com‐ parable to the available tensile data for this alloy.

A low hardening rate observed in regime 3 of these alloys can be solely attributed to the growth of existing precipitates. No further increase in the precipitate density occurs in this regime leading to a stable ductility. The final brittle regime (regime 4) with an increasing hardening rate and a decreasing ductility is observed only in 6061-T6 alloy (from CRD A-2 tubes of HFBR) at very high fluences. Although this alloy is the same as 6061-T6 alloy and irradiated at similar temperatures, a difference in behavior is observed due to irradiation at very high TFR, as explained in Section 4.4.

#### **4.2. Tensile behavior of 5***xxx* **series alloys**

The differences in irradiation hardening trends in all four regimes of 5*xxx* and 6*xxx* series alloys are depicted schematically in **Figure 2**. Both 5052-O and 5154-O alloys show similar irradiation hardening and embrittlement behavior (**Figure 1**) due to similar alloy micro‐ structure and irradiation conditions. The unirradiated strength values of 5052-O and 5154-O alloys are lower, and ductility is higher than 6061-T6 alloy (**Figure 1**) due to the absence of Mg2Si precipitates before irradiation. In the rapid hardening regime, the magnitude of irra‐ diation hardening and embrittlement in 5*xxx* series alloys is observed to be much higher than 6*xxx* series, because both dislocation hardening (due to displacement damage) and pre‐ cipitation hardening (due to the formation of Mg2Si precipitates from transmutation Si) oc‐ cur simultaneously in 5*xxx* alloys (see Section 3.1).

The contribution of both mechanisms continues in the transition regime until dislocation damage reaches a saturation value (at <2 × 1026 n/m2 of fast fluence or <4 × 1026 n/m2 of ther‐ mal fluence). Simultaneously, a saturation in the density of precipitates is expected to occur in this regime, leading to the formation of no new Mg2Si precipitates.

0.13 <sup>Δ</sup> ln λ

1/2 Δ ,

 ss

s

s

parable to the available tensile data for this alloy.

very high TFR, as explained in Section 4.4.

**4.2. Tensile behavior of 5***xxx* **series alloys**

cur simultaneously in 5*xxx* alloys (see Section 3.1).

damage reaches a saturation value (at <2 × 1026 n/m2

equation [20]:

402 Radiation Effects in Materials

*Gb r b*

Assuming a saturation density of ~ 6 × 1014 m−2 in Al alloys (same as in steel), a rough esti‐ mate of the total contribution of dislocation hardening can be made using the following

where *σ* is the strength of the material after introducing dislocation structure, *σo* is the intrin‐ sic strength of the material with low dislocation density, *G* is shear modulus (=2.648 × 104 MPa), *b* is Burgers vector (0.286 nm) and *ρ* is dislocation density. The *G* and *b* values of pure Al taken from reference [21] are used here. Substituting *ρ* = 6 × 1014 m−2 into Eq. (3) gives Δ*σ* = ~186 MPa. The magnitude of hardening observed at the beginning of 6061-O and 6061-T6 alloys is in good agreement with this value. An estimation of the total irradiation hardening using the microstructure data (*r* = 4 nm, λ = 4 nm) of CRDF A-2 alloy at 42 × 1026 n/m2 ther‐ mal fluence taken from [5] resulted in an estimated yield strength of 605 MPa, which is com‐

A low hardening rate observed in regime 3 of these alloys can be solely attributed to the growth of existing precipitates. No further increase in the precipitate density occurs in this regime leading to a stable ductility. The final brittle regime (regime 4) with an increasing hardening rate and a decreasing ductility is observed only in 6061-T6 alloy (from CRD A-2 tubes of HFBR) at very high fluences. Although this alloy is the same as 6061-T6 alloy and irradiated at similar temperatures, a difference in behavior is observed due to irradiation at

The differences in irradiation hardening trends in all four regimes of 5*xxx* and 6*xxx* series alloys are depicted schematically in **Figure 2**. Both 5052-O and 5154-O alloys show similar irradiation hardening and embrittlement behavior (**Figure 1**) due to similar alloy micro‐ structure and irradiation conditions. The unirradiated strength values of 5052-O and 5154-O alloys are lower, and ductility is higher than 6061-T6 alloy (**Figure 1**) due to the absence of Mg2Si precipitates before irradiation. In the rapid hardening regime, the magnitude of irra‐ diation hardening and embrittlement in 5*xxx* series alloys is observed to be much higher than 6*xxx* series, because both dislocation hardening (due to displacement damage) and pre‐ cipitation hardening (due to the formation of Mg2Si precipitates from transmutation Si) oc‐

The contribution of both mechanisms continues in the transition regime until dislocation

of fast fluence or <4 × 1026 n/m2

of ther‐

 r

= (2)

*dis* =- =*<sup>o</sup> Gb* (3)

With further irradiation, the hardening continues with a decreasing rate as Mg2Si precipi‐ tates continue to grow until all the Mg is pulled out from the Al solid solution in the final slow hardening regime. Based on the stoichiometric analysis, production of 0.58 wt.% trans‐ mutation Si will consume 1% Mg in the alloy. That means all the Mg in 5154-0 alloy is con‐ sumed at ~1.85% transmutation Si (~8.66 × 1026 n/m2 of thermal fluence) and in 5052-O alloy at ~1.27% transmutation Si (~5.95 × 1026 n/m2 of thermal fluence).

With continued irradiation, the newly formed transmutation-produced Si either decorates existing precipitates (like in 6*xxx* series) or associates with voids, which are expected to form only at very high fast fluences in 5*xxx* alloys as explained in Section 3.2. The higher harden‐ ing rate observed in regime 3 of 5*xxx* alloys can be attributed to the higher density and vol‐ ume fraction of Mg2Si precipitates than that observed in 6*xxx* series alloys. A stable ductility is observed in this stage similar to 6*xxx* series alloys due to no further increase in precipita‐ tion density. In fact, a small decrease in particle density may occur in this regime due to par‐ ticle coalescence during their growth. The opposite effects of a small decrease in particle density and a slow hardening due to precipitate growth on ductility could be compensating each other, leading to a plateau in the ductility behavior in regime 3. No brittle regime is observed in the available data of 5*xxx* alloys until a thermal fluence of 31 × 1026 n/m2 .

Similar to 6*xxx* alloys, a rough estimation of the total irradiation hardening contribution is performed for 5052-O and 5154-O alloys using Eq. (2) for the dislocation-hardening contri‐ bution and Eq. (3) for contribution from precipitation hardening. An average precipitate size of 10 nm and linear planar spacing between precipitates of 10 nm is used from the micro‐ structure data presented in [3] on 5052-O alloy irradiated to a thermal fluence of 31 × 1026 n/m2 . This resulted in a total yield strength of 547 MPa (including unirradiated yield strength value of 85 MPa) for 5052-O, matching very well with the data presented in **Figure 1 (a)**. Similar analysis for 5154-O at a thermal fluence of 9.81 × 1026 n/m2 , with an average precipitate size of 8 nm and linear planar spacing between precipitates of 20 nm (obtained from **Figure 5**) [13] resulted in a total yield strength of 410 MPa, again matching with the trends seen in **Figure 1 (a)**.

#### **4.3. Fracture toughness behavior of 5***xxx* **and 6***xxx* **series alloys**

In this section, first the fracture toughness data from HFR SURP is plotted against the avail‐ able literature data on highly irradiated Al alloys. The evolution of fracture toughness be‐ havior of 5*xxx* and 6*xxx* alloys during neutron irradiation is discussed, and a connection is established between tensile and fracture toughness behavior in various regimes observed in the previous section.

#### *4.3.1. Literature fracture toughness data of 5xxx series Al alloys*

No additional data on fracture toughness properties of 5*xxx* Al alloys was found in the liter‐ ature except the data from 5154-O alloy from HFR SURP [13]. The fracture toughness prop‐ erties of the HFR surveillance specimens are periodically measured to assess and predict the hotspot behavior of HFR vessel in comparison with the vessel's fracture toughness design limit of 6 (MPa)·m1/2 [22]. Fracture toughness properties of surveillance specimens from the current HFR vessel tested until 2010 are plotted as a function of thermal neutron fluence in **Figure 3 (a)**. The thermal neutron fluence is taken as an indicator for the measure of irradia‐ tion damage in HFR vessel material, because thermal neutrons are the major cause of dam‐ age by producing Si through transmutation, as explained in the previous sections. **Figure 3 (b)** shows the relation between the thermal fluence and transmutation Si in fracture tough‐ ness specimens tested in SURP.

**Figure 3.** (a) HFR SURP fracture toughness data as a function of thermal neutron fluence. (b) Transmutation Si values as a function of neutron thermal fluence of fracture toughness samples tested in HFR SURP program. Note that the projected hotspot thermal fluence and Si content are based on the assumption that the irradiation conditions at the HFR hotspot are kept unchanged as they are in 2015.

*4.3.2. Literature fracture toughness data of 6xxx series Al alloys in comparison with HFR SURP data*

Only limited data was published on fracture toughness properties of irradiated Al alloys [5, 8, 10]. The most relevant data for the HFR (irradiation temperatures < 373 K) is plotted in **Figure 4** in comparison with HFR SURP data. Data from 6061-T6 alloy irradiated at < 373 K in the High Flux Isotope Reactor (HFIR) in Oak Ridge National Laboratory, USA matches quite well with the HFR SURP data. As it can be seen from **Figure 4**, there is one high fluence data point published by Weeks et al. [5] beyond the current surveillance data of the HFR vessel. This data is from the CRDF A-2 tubes of the HFBR in Brookhaven National Laboratory, USA, produced from 6061-T6 alloy, irradiated at 338 K up to a thermal fluence of 42 × 1026 n/m2 . The corresponding fast fluence of this data point is 2 × 1026 n/m2 , which gives a high TFR of 21, compared to the HFR hotspot TFR value of maximum 1.4. The total measured Si at this fluence was found to be ~8 wt.%, including 0.6% of initial Si content. The reported thermal fluence and Si content of this data point are approximately two times the estimated thermal fluence (~20 × 1026 n/m2 ) and Si (~4.3 %) content of the HFR hotspot by the end of 2025. Note that this data is from the same material and at the same irradiation conditions for which the tensile data at very high thermal fluences (~42 × 1026 n/m2 ) is also available (see **Figure 1**).

*4.3.1. Literature fracture toughness data of 5xxx series Al alloys*

ness specimens tested in SURP.

404 Radiation Effects in Materials

HFR hotspot are kept unchanged as they are in 2015.

No additional data on fracture toughness properties of 5*xxx* Al alloys was found in the liter‐ ature except the data from 5154-O alloy from HFR SURP [13]. The fracture toughness prop‐ erties of the HFR surveillance specimens are periodically measured to assess and predict the hotspot behavior of HFR vessel in comparison with the vessel's fracture toughness design limit of 6 (MPa)·m1/2 [22]. Fracture toughness properties of surveillance specimens from the current HFR vessel tested until 2010 are plotted as a function of thermal neutron fluence in **Figure 3 (a)**. The thermal neutron fluence is taken as an indicator for the measure of irradia‐ tion damage in HFR vessel material, because thermal neutrons are the major cause of dam‐ age by producing Si through transmutation, as explained in the previous sections. **Figure 3 (b)** shows the relation between the thermal fluence and transmutation Si in fracture tough‐

**Figure 3.** (a) HFR SURP fracture toughness data as a function of thermal neutron fluence. (b) Transmutation Si values as a function of neutron thermal fluence of fracture toughness samples tested in HFR SURP program. Note that the projected hotspot thermal fluence and Si content are based on the assumption that the irradiation conditions at the

*4.3.2. Literature fracture toughness data of 6xxx series Al alloys in comparison with HFR SURP data* Only limited data was published on fracture toughness properties of irradiated Al alloys [5, 8, 10]. The most relevant data for the HFR (irradiation temperatures < 373 K) is plotted in

**Figure 4.** HFR SURP fracture toughness data in comparison with literature data [5,10].

#### *4.3.3. Fracture toughness behavior of 5xxx and 6xxx series alloys*

In the rapid hardening regime (regime 1), the fracture toughness value drops rapidly in line with the observed hardening and ductility behavior of 5*xxx* alloys (**Figure 3** vs. **Figure 1 (c)**). This is because the hardening-induced embrittlement causes the decrease in both ductility and fracture toughness properties. In this regime, the magnitude of the drop in 5*xxx* alloys is high compared to 6*xxx* alloys, because both dislocation damage and precipitation damage mecha‐ nisms are active in 5*xxx* alloys, whereas only dislocation damage dominates in 6*xxx* alloys. This explains the sharp drop observed in fracture toughness of 5154-O alloy at the onset of irradiation compared to a shallow decrease in fracture toughness properties of 6061-T6 alloy in this regime (see **Figure 4**).

As the irradiation continues, both the dislocation density and the Mg2Si precipitate density evolve toward a saturation limit describing the slow decrease of fracture toughness toward a plateau in the transition regime. Transmission electron microscopy results of precipitate microstructure reported in [13] are shown in **Figures 5** and **6**. From these results it can be seen that the saturation density is achieved at ~3 × 1026 n/m2 of thermal fluence for 5154-O alloy of HFR vessel. Note that these pictures were taken using a "JEOL JEM-1200ex STEM/TEM" machine operating at 120 keV, located in JGL laboratory at NRG.

After that, the fracture toughness of 5154-O reaches a plateau at a thermal fluence of ~ 4 × 1026 n/m2 after which no further increase in dislocation and precipitate density is expected (**Figure 3**). In fact, a small decrease in particle density may occur later in this regime due to particle coalescence during their growth. The opposite effects of a small decrease in particle density and a slow hardening due to precipitate growth on embrittlement could be compen‐ sating each other leading to a plateau in the fracture toughness behavior (similar to ductility) in regime 3. The behavior of 5154-O alloy is expected to be similar to 6*xxx* series alloys in this regime at similar irradiation conditions. This is because within this regime the irradiation hardening in both alloy types occurs primarily due to the growth of existing precipitates by newly produced transmutation Si. A close agreement between the fracture toughness data of 6061-T6 alloy irradiated in HFIR at <373K and HFR SURP data in the plateau regime (see **Figure 4**) confirms this theory.

**Figure 5.** TEM images showing the evolution of precipitate size with thermal fluence in 5154-O Al alloy (HFR SURP specimens). (a) 2.73x1026 n/m2, 0.68 wt.% Si, (b) 3.76x1026 n/m2, 0.88 wt.% Si and (c) 9.81x1026 n/m2, 2.21 wt.% Si. (Photograph courtesy N.V. Luzginova et. al. [13])

This is because the hardening-induced embrittlement causes the decrease in both ductility and fracture toughness properties. In this regime, the magnitude of the drop in 5*xxx* alloys is high compared to 6*xxx* alloys, because both dislocation damage and precipitation damage mecha‐ nisms are active in 5*xxx* alloys, whereas only dislocation damage dominates in 6*xxx* alloys. This explains the sharp drop observed in fracture toughness of 5154-O alloy at the onset of irradiation compared to a shallow decrease in fracture toughness properties of 6061-T6 alloy

As the irradiation continues, both the dislocation density and the Mg2Si precipitate density evolve toward a saturation limit describing the slow decrease of fracture toughness toward a plateau in the transition regime. Transmission electron microscopy results of precipitate microstructure reported in [13] are shown in **Figures 5** and **6**. From these results it can be seen that the saturation density is achieved at ~3 × 1026 n/m2 of thermal fluence for 5154-O alloy of HFR vessel. Note that these pictures were taken using a "JEOL JEM-1200ex STEM/TEM"

After that, the fracture toughness of 5154-O reaches a plateau at a thermal fluence of ~ 4 ×

(**Figure 3**). In fact, a small decrease in particle density may occur later in this regime due to particle coalescence during their growth. The opposite effects of a small decrease in particle density and a slow hardening due to precipitate growth on embrittlement could be compen‐ sating each other leading to a plateau in the fracture toughness behavior (similar to ductility) in regime 3. The behavior of 5154-O alloy is expected to be similar to 6*xxx* series alloys in this regime at similar irradiation conditions. This is because within this regime the irradiation hardening in both alloy types occurs primarily due to the growth of existing precipitates by newly produced transmutation Si. A close agreement between the fracture toughness data of 6061-T6 alloy irradiated in HFIR at <373K and HFR SURP data in the plateau regime (see **Figure**

**Figure 5.** TEM images showing the evolution of precipitate size with thermal fluence in 5154-O Al alloy (HFR SURP specimens). (a) 2.73x1026 n/m2, 0.68 wt.% Si, (b) 3.76x1026 n/m2, 0.88 wt.% Si and (c) 9.81x1026 n/m2, 2.21 wt.% Si.

after which no further increase in dislocation and precipitate density is expected

machine operating at 120 keV, located in JGL laboratory at NRG.

in this regime (see **Figure 4**).

406 Radiation Effects in Materials

1026 n/m2

**4**) confirms this theory.

(Photograph courtesy N.V. Luzginova et. al. [13])

**Figure 6.** TEM images showing the evolution of precipitate microstructure with thermal fluence in 5154-O Al alloy (HFR SURP specimens). (a) Unirradiated, 0.04 wt.% Si, (b) 2.73 × 1026 n/m2 , 0.68 wt.% Si, (c) 3.76 × 1026 n/m2 , 0.88 wt.% Si, and (d) 9.81 × 1026 n/m2 , 2.21 wt.% Si. (Photograph courtesy N.V. Luzginova et al. [13].)

It is important to understand how long the plateau in the fracture toughness (or regime 3) will continue. This depends on the location of the precipitation of the transmutation Si. As already mentioned in Section 3.2.2, further increase in Si production to high values can lead to Si precipitation at grain boundaries. Fracture toughness value drops when the precipitation of Si at the grain boundaries cumulates to an extent that the dominant deformation and fracture mechanisms shift from the bulk microstructure to the grain boundaries. A heavy discontinuous precipitation observed at the grain boundary in 5052-O alloy at a thermal fluence of 31 × 1026 n/m2 [3] has resulted in no substantial effects on ductility of this alloy. This suggests that the nature of fracture at these high fluence is still controlled by bulk deformation mechanisms (instead of mechanisms controlled by grain boundaries). Due to the similarity in 5154-O and 5052-O alloys (and irradiation conditions), the ductility and fracture toughness properties of the 5154-O alloy are also expected to show a plateau until such high fluences. Indeed, the observation of significant amount of micron-scale dimples on the fracture surface of 5154-O alloy irradiated to a thermal fluence of 9.81× 1026 n/m2 (**Figure 7**) proves that similar behavior can be expected from the 5154-O alloy [13].

**Figure 7.** Fracture details of 5154-O alloy with a crack-tip thermal fluence of 9.81 × 1026 n/m2 . Figure shows fracture surface characterized by dominant microdimples and some cleavage facets [13].

CRDF A-2 tubes of HFBR produced from 6061-T6 alloy have shown a fracture toughness value of ~8 (MPa)·m1/2 after irradiation to a much higher thermal neutron fluence of ~42 × 1026 n/m2 at 338 K (see **Figure 4**). The decrease in fracture toughness from an unirradiated value of 21.75 (MPa)·m1/2 for this alloy is primarily attributed to the following: (i) formation of very fine (~8 nm) Si-rich precipitates in the grains due to high TFR of 21 (as explained in Section 4.4) and (ii) large silicon flakes occupying about one-fifth of the grain boundary area at this high transmutation-produced Si content of 8 wt.% [5]. Fracture surface of this alloy revealed substantial intergranular separation with some residual ductility indicating that the contribu‐ tion of grain boundary fracture mechanisms is increased at such high fluence values to enter into the brittle regime (regime 4).

From the above discussion, no differences in the evolution of irradiation damage at high fluences (in regime 3 and 4) are expected between 5*xxx* and 6*xxx* alloys due to differences in their initial chemical composition and microstructure. Moreover, the difference in the TFR of HFR SURP data and literature data is conservative as explained in the next section. This allows the use of published fracture toughness literature data of 6061-T6 Al alloys to predict the fracture toughness behavior of 5154-O Al alloy of HFR vessel at high fluences.

#### **4.4. Effect of thermal-to-fast flux ratio (TFR)**

It is known from the literature that a high difference in TFR can have substantial effect on irradiation hardening and embrittlement behavior of the same material [5]. It was highlighted in [3] that both the thermal and fast neutrons play independent and important roles leading to microstructural damage and corresponding property changes. A very high TFR, ranging from 80 to 500, could explain the observed craze-cracking in AG3-NET alloy (Al–3% Mg) beam tubes in the Reactor Haut Flux (RHF) at Grenoble [23]. Lijbrink et al. [12] pointed out that fast neutron flux reduces the effectiveness of the Si precipitation hardening process. A possible explanation for this behavior (as given in [5, 12]) is as follows. Fast flux has two opposite effects on precipitation:

**i.** The kinetic energy supplied by fast flux temporarily increases the solubility limit of Si in the matrix and opposes the condensation requirements for the precipitation.

**ii.** Local energy needed for jumping the nucleation barrier can be readily supplied by the fast flux.

However, the fast flux can be destructive when a freshly formed nucleus is hit by fast neutron collision. That means, at equal thermal fluence values, Si precipitation hardening is more effective at higher TFR. This leads to finer precipitate distribution, causing higher irradiation hardening, lower ductility, and eventually lower fracture toughness values at higher TFR. Indeed, the higher hardening rate observed in 6061-T6 alloy from CRDF A-2 tubes of HFBR irradiated at TFR of 21 compared to the similar 6061-T6 alloy irradiated in HFIR at TFR of 1.7 explains this behavior (**Figure 1 (b)**). Consequently, the high fluence data point from CRDF A-2 of HFBR at TFR = 21 (>>0.8–1.4 for HFR hotspot), shown in **Figure 4**, is likely to give a conservative estimation of the fracture toughness value under HFR conditions.
