Linear complexions enable unprecedented ductility retention in neutron irradiated ferritic steel

Linear complexions enable unprecedented ductility retention in neutron irradiated ferritic steel | Research Square window.SnipcartSettings = { analytics: { enabled: false } }; (function() { var accessVector = localStorage.getItem('access_vector') || ''; window.dataLayer = window.dataLayer || []; if (accessVector) { window.dataLayer.push({ user: { profile: { profileInfo: { snid: accessVector } } } }); } })(); (function(w,d,s,l,i){w[l]=w[l]||[];w[l].push({'gtm.start':new Date().getTime(),event:'gtm.js'});var f=d.getElementsByTagName(s)[0],j=d.createElement(s),dl=l!='dataLayer'?'&l='+l:'';j.async=true;j.src='https://www.googletagmanager.com/gtm.js?id='+i+dl;f.parentNode.insertBefore(j,f);})(window,document,'script','dataLayer','GTM-K279D39R'); Browse Preprints In Review Journals COVID-19 Preprints AJE Video Bytes Research Tools Research Promotion AJE Professional Editing AJE Rubriq About Preprint Platform In Review Editorial Policies Our Team Advisory Board Help Center Sign In Submit a Preprint Cite Share Download PDF Article Linear complexions enable unprecedented ductility retention in neutron irradiated ferritic steel Janelle Wharry, Arya Chatterjee, Soumita Mondal, Yu Lu, Yaqiao Wu This is a preprint; it has not been peer reviewed by a journal. https://doi.org/ 10.21203/rs.3.rs-7419483/v1 This work is licensed under a CC BY 4.0 License Status: Under Review Version 1 posted You are reading this latest preprint version Abstract Herein, the operando formation of Si-enriched linear complexions during neutron enables Grade 91 ferritic steel to overcome the strength–ductility tradeoff, one of the most critical life-limiting challenges facing nuclear structural alloys. Linear complexions are a distinct yet confined chemical and structural state at a dislocation, which are rarely reported in engineering alloys. Ferritic steels are amongst the most ubiquitous engineering alloys for current and future nuclear components, but they are susceptible to irradiation hardening and embrittlement. Here, exceptional ductility retention exceeding 90% of pre-irradiation levels is obtained in Grade 91 synthesized using powder metallurgy with hot isostatic pressing (PM-HIP). Powder processing artifacts promote a high density of screw dislocation arrays, on which β-FeSi 2 linear complexions form due to Si segregation during irradiation. Screw dislocation dipoles undergo pinning and unpinning on linear complexions, resulting in extended yielding and excellent ductility retention post-irradiation. These findings represent a significant advancement toward design of alloys and manufacturing processes that can autonomously self-regulate their microstructural resilience in-operando during irradiation, enabling exceptional ductility rather than embrittlement. Physical sciences/Materials science/Structural materials/Metals and alloys Physical sciences/Materials science/Structural materials/Mechanical properties Figures Figure 1 Figure 2 Figure 3 Figure 4 Figure 5 Figure 6 Figure 7 Introduction Strength and ductility are the key properties that influence the broad utilization of any structural material across varied applications and industries 1 , 2 , but these properties are often mutually exclusive 3 . The strength and ductility of metallic alloys, including those used for nuclear applications, depend on the presence of crystalline defects and their mobility under mechanical loading. The nature of plasticity in a metallic alloy is controlled by the interaction between dislocations (i.e., one-dimensional line defects) and all other types of defects including stacking faults 4 , grain boundaries 5 , precipitates 6 , solute atoms 7 , 8 , and other dislocations. Generally, as dislocations interact with defects, those defects can present strong obstacles that hinder the motion of dislocations, resulting in strengthening of the material. But this often comes at the sacrifice of the ability for the material to deform plastically, leading to a reduction in ductility. This strength–ductility tradeoff becomes even more exacerbated in metallic alloys exposed to irradiation, such as during service in a nuclear reactor environment. Irradiation generates a supersaturation of point defects, which can evolve into extended defects such as clusters, dislocation loops, stacking faults, voids, solute segregation at the aforementioned features, precipitation, and twinning and phase transformations 9 . These irradiation-induced microstructures and microchemical gradients tend to hinder dislocation motion, resulting in an increase in strength and decrease in ductility. This phenomenon of irradiation hardening and embrittlement poses a threat of brittle fracture in irradiated structural alloys, consequently compromising the safe operation of reactors. Recently, numerous studies have sought to use defect engineering to intentionally introduce defects that can potentially help overcome the strength–ductility tradeoff in metallic alloys. One such approach has been to extend plasticity by making twinning and mechanical phase transformations favorable via alloy design of twinning-induced plasticity (TWIP) and transformation-induced plasticity (TRIP) steels 2 , 10 . By obstructing dislocation motion, twins enhance strength, whilst the existence of multiple twins facilitates dislocation mobility through twin reorientation 2 . Similarly, phase transformations during deformation alter dislocation behavior in these new phases, often enhancing dislocation mobility 10 . In other cases, irradiation-induced defects such as voids and dislocation loops can also promote twinning and phase transformation in non-TWIP/TWIP alloys 11 – 15 . Another approach to defect engineering is to take advantage of solute segregation, which is prevalent in irradiated materials. Chemical segregation is generally thought to hinder dislocation motion, thus contributing to hardening and loss of ductility 16 , 17 . However, if the segregation can be tailored such that it leads to the formation of a complexion, improvements in mechanical properties may be achieved 18 . A linear complexion is a distinctly identifiable, but confined, chemical and structural state at dislocations 19 . Complexions can eventually give rise to stable 20 or metastable 21 , 22 phase formation. But if the extent of solute segregation to dislocation lines is insufficient to form stable precipitate growth and coalescence along those lines, linear arrays of nanoscale “precipitates” can form along the dislocation, creating what are known as linear complexions 19 – 23 . These linear complexions can significantly influence strength and ductility and strength through dislocation pinning and unpinning 21 , 24 , 25 . The formation of linear complexions at dislocations has thus far primarily been predicted through atomistic simulations in model binary alloys such as Fe-Ni, Cu-Zr, Al-Cu, Ni-Al, and Al-Zr 20 , 22 – 26 . Only a few studies have identified linear complexion formation experimentally by engineered thermo-mechanical treatments. For example, complexions can form in normalized Fe-9at.%Mn alloy through additional rolling to impart a high dislocation density, followed by extended tempering treatments allowing for Mn diffusion in Fe 19 , 21 , 27 . Nevertheless, nearly all reports of linear complexions are in model binary alloys 19 – 30 . Harnessing solute segregation to form linear complexions at dislocations, instead of forming clusters or precipitates, could thus remarkably improve ductility in an irradiation-hardened alloy. This pioneering study presents the novel formation of linear complexions at dislocations in an engineering alloy under service-relevant neutron irradiation conditions. Even more notably, these complexions promote excellent retention of ductility in irradiation-hardened materials, overcoming the longstanding strength–ductility tradeoff. This work focuses on Grade 91 (G91) ferritic steel, used in current and advanced nuclear fission reactor structural components, manufactured using conventional casting and by powder metallurgy with hot isostatic pressing (PM-HIP). Materials are neutron irradiated to 1 and 3 displacements per atom (dpa) at 400ºC. The PM-HIP 1 dpa G91 forms linear complexions of the β-FeSi 2 phase along Si-segregated dislocation lines, enabling high retained ductility despite undergoing irradiation hardening. Coupling transmission electron microscopy (TEM) with atom probe tomography (APT) provides evidence of the structure and chemistry of the linear complexions. Finally, the manuscript discusses the plausible complexion formation and deformation mechanisms that enable commercial G91 ferritic steel to be amongst the first irradiated engineering alloys to overcome the strength-ductility tradeoff through complexion engineering. Results Mechanical behavior of unirradiated and irradiated G91 steel. Engineering stress-engineering strain curves from uniaxial tensile testing of the G91 steels are provided in Fig. 1 a and summarized in Table 1 . Tensile tests reveal the general trend of irradiation hardening, as yield stress increases with neutron irradiation. Prior to irradiation, cast and PM-HIP G91 steels show similar yield strengths (YS) of 403 MPa and 447 MPa, ultimate tensile strengths (UTS) of 606 and 639 MPa, and uniform elongations (UE) of 10%, respectively. Upon 1 dpa neutron irradiation, PM-HIP exhibits higher hardening than cast (YS of 664 MPa compared to 628 MPa), similar UTS (737 MPa compared to 734 MPa), and retains higher ductility (UE of 9.1% as compared to 5.6%). At 3 dpa, the cast alloy exhibits higher hardening and strengthening than the PM-HIP (YS 765 MPa and UTS 819 MPa for PM-HIP; YS 801 MPa and UTS 867 MPa for cast), whilst both show a significant reduction in ductility, although PM-HIP continues to retain greater ductility than cast (UE of 4.4% and 2.3% for PM-HIP and cast, respectively). These superior irradiation hardening and embrittlement behaviors of the PM-HIP alloy as compared to its cast counterpart has also been observed in other ferritic steels and Ni-based alloys 31 , 32 Fractography of the broken tensile bars, Fig. 1 b, reveals void-dominant dimpled fracture in unirradiated G91 steels, with some quasi-cleavage regions on the fracture surface of the unirradiated cast specimen. This quasi-cleavage fracture region becomes more prominent in the cast specimen after 1 dpa, whereas the PM-HIP retains dimpled fracture surfaces after 1 dpa, indicative of excellent retained ductility close to its pre-irradiated state. In both 3 dpa irradiated specimens, angular faceted features are present along with evidence of larger voids, consistent with the lower ductility. The mechanical properties of PM-HIP G91 steel after 1 dpa are noteworthy, as the alloy retains > 90% of its pre-irradiation ductility, despite significant irradiation-induced strengthening, Fig. 1 c. A key driver for maintaining this exceptional ductility is the extended yielding phenomenon observed in 1 dpa irradiated PM-HIP G91, as depicted within the inset of Fig. 1 a, followed by gradual strain hardening. This extended yielding phenomenon is identified by the accumulation of ~ 2% strain with overall negligible increase in bulk stress value, occurring between yielding and work hardening. These discontinuous yielding phenomena have been associated with various mechanisms: dislocation-solute interactions, such as Cottrell atmospheres 33 when upper and lower yield points are observed; Lüders band formation in cases of strain localization; and dynamic strain aging effects, such as Portevin-Le Chatelier (PLC) behavior for the case of serrated or jerky flow 34 . Dislocation-solute interactions have primarily considered the effects of interstitial solute atoms (e.g., C and N) in Fe-based binary systems 33 , ferritic-martensitic alloys 35 , and austenitic Fe-Mn-C TWIP steels 36 . Similarly, under irradiation, when the rate of solute atom diffusion is slower than the dislocation glide velocity, Lüders band phenomena can be observed through the formation of cleared channels in the microstructure 37 . By contrast, when solute diffusion rates are comparable to dislocation glide velocity, the PLC effect occurs through a repeated locking-unlocking mechanism between solutes and dislocations, and by propagating intermittent plasticity throughout the specimen 38 . Table 1 Mechanical properties of investigated irradiated G91 steels. Material YS (MPa) UTS (MPa) Uniform elongation (%) Cast unirradiated 402.8 606.0 9.88 PM-HIP unirradiated 447.2 639.3 10.00 Cast 1 dpa 627.8 733.8 5.62 PM-HIP 1 dpa 663.7 736.9 9.12 Cast 3 dpa 800.9 867.0 2.32 PM-HIP 3 dpa 764.8 819.2 4.42 Characterization of microstructure to correlate mechanical properties . Transmission electron microscopic (TEM) analyses show that the initial microstructure of both cast and PM-HIP G91 steel is comprised of lath martensite with larger M 23 C 6 (M = Cr/Mo) type carbides and nanoscale microalloyed MX type precipitates (M = V, X = C/N), Fig. 2 (a, b ). Such a microstructure is typical of 9Cr-1Mo F/M steels, consistently reported in numerous earlier studies 43 , 44 . Following irradiation, no cavities are present at ~ 1 dpa, while only a negligible population of cavities is found at ~ 3 dpa, Fig. 2 (c, d) . Cast G91 has ~ 5 nm diameter cavities at a number density of 3.30×10²¹ m − 3 after 3 dpa, while the PM-HIP exhibits ~ 4 nm diameter cavities at a number density of 2.13×10²¹ m − 3 . These cavities may result in only ∼0.001% volumetric swelling. The extent of cavity formation observed in this study is comparable to earlier work on Grade 91 steel from Tan, et al. 45 , wherein 4.36 dpa neutron irradiation at 469ºC results in an average cavity diameter of 3.0 ± 1.1 nm (up to ∼6.8 nm) with a number density of (3.9 ± 0.6) × 10 21 m − 3 , corresponding to ~ 0.005% swelling. F/M steels are inherently resistant to void swelling, even at higher irradiation doses owing to their high density of martensite lath boundaries and extensive dislocation networks, as can be observed in Fig. 2 (a, b) . These lath boundaries and dislocation networks act as efficient sinks for recombination of point defects, thereby delaying the onset of void swelling 46 . Irradiation-induced dislocation loops are populous in irradiated G91, Fig. 2 (e-h) . At 1 dpa, the average loop diameter is 27.8 ± 5.1 nm for cast and 26.3 ± 3.4 nm for PM-HIP, with corresponding number densities of 4.8 × 10²¹ m − 3 and 4.4 × 10²¹ m − 3 , respectively, Table 2 . Loops grow significantly by 3 dpa, reaching 43.5 ± 3.7 nm and 38.3 ± 2.8 nm, whilst their number densities decrease marginally to 4.07 × 10²¹ m − 3 and 3.7 × 10²¹ m − 3 for cast and PM-HIP steels, respectively, Table 2 . Loop sizes and densities are in agreement with previous reports on ~ 350–500ºC neutron irradiated F/M steels, in which loop diameters range 4.4–6.1 nm at a density of 6.8–7.8 × 10 21 m − 3 for doses 0.7–1.1 dpa, and 10.6–45.5 nm at a density of 0.8–4.8 × 10 21 m − 3 for doses 3.9–4.1 dpa 47 , 48 . Dislocation loops in 1 dpa cast and PM-HIP G91 are preferentially clustered near or along pre-existing dislocation lines and networks, Fig. 2 (e, f). By 3 dpa, loops appear more homogeneously distributed, though still exhibit some preferential location around martensite lath boundaries and subgrain boundaries, Fig. 2 (g, h) . Previous studies on low dose (generally < 1 dpa) irradiated F/M steels also report similar heterogeneous distributions of loops near dislocation lines and networks 48 , indicating that dislocations may be sites for nucleation and trapping of loops. At higher dose levels, more vacancies and self-interstitial atoms (SIAs) are generated. Dislocation loops act as sinks for these point defects, absorbing them from the surrounding matrix 49 . Moreover, continuous absorption of SIAs leads to the growth of existing interstitial dislocation loops and can coalesce to other loops or dislocation network 50 , thereby increasing the loops size and becoming more uniformly distributed throughout the microstructure as observed at 3 dpa. The presence of loops along dislocation lines and networks in the 1 dpa specimens is similar to the irradiation defect–dislocation interaction phenomenon often referred to as ‘rafts’ 37 , wherein loops adhere to dislocation lines. Rafts form when irradiation-induced self-interstitial atom clusters migrate rapidly in a dislocation stress field, resulting in self-interstitial loops forming close to dislocations lines. These interstitial loops effectively pin dislocations and reduce their mobility. Only with additional stress, these locked dislocations can become mobile again. One may consider attributing the extended yielding phenomenon in 1 dpa PM-HIP G91 to dislocation–raft locking–unlocking, but this is unlikely. First, even though the loop–dislocation distributions are similar in both cast and PM-HIP 1 dpa microstructures, the cast steel does not exhibit extended yielding. Second, rafts generally lead to strain localization by generating dislocation channels, which are absent in the irradiated-then-strained deformation microstructures collected near the fracture surface of all investigated specimens, Fig. S1 (supplementary file) . Characterization of segregation . Irradiation-induced segregation of solute elements are characterized with atom probe tomography (APT). Reconstructed three dimensional (3D) atom maps obtained from the APT tips of the irradiated G91 specimens are depicted in Fig. 3 (a-d) . Discrete spherical solute clusters form in the cast specimen at 1 and 3 dpa, and in the PM-HIP specimen only at 3 dpa, Fig. 3 (a, c, d) ; by contrast, in the PM-HIP specimen at 1 dpa, solute segregation occurs primarily along dislocation lines instead of nucleating discrete clusters, Fig. 3 b. The primary solutes that form clusters are Ni, Mn, Si, and Cu, as listed in Table 3 . However, the dominant solute that decorates dislocation lines in the 1 dpa PM-HIP G91 is Si with sparse Ni clusters, Fig. 3 b. Solute clustering is consistent with previous reports of irradiation-induced Ni-Mn-Si-rich (sometimes identified as stoichiometrically G-phase) and Cu-rich nanoclusters which form at number densities ~ 10 23 m − 3 over irradiation doses ~ 1–10 dpa at temperatures ~ 300–600ºC in F/M steels including Grade 91 type steels 47 , 48 , 51 , 52 . These clusters have previously been linked to irradiation hardening and are thus consistent with the observed uniaxial tensile stress-strain behavior of the cast 1 and 3 dpa and PM-HIP 3 dpa specimens herein. The reader is referred to previous studies 47 , 48 , 53 , 54 for details of nanocluster formation mechanisms and their implications on irradiation hardening. Here, this study will instead aim to explain the nature of solute segregation to dislocations in the 1 dpa PM-HIP specimen, and its implications on extended yielding. To characterize the irradiation-induced segregation to dislocations for the 1 dpa PM-HIP specimen, composition profiles are taken along and across the array of piled-up dislocation lines observed in the APT reconstruction Fig. 3 b. Although Si enrichment is evident through the entirety of the dislocation pile-up array, the composition of Si fluctuates periodically along the dislocation lines; 1-D composition profiles reveal the period of Si fluctuation is ~ 4.7 ± 1.2 nm, Fig. 4 a. The Si concentration appears anti-correlated with either the Fe or Cr concentration, suggesting a substitutional solute exchange formation mechanism (profiles 1–3, Fig. 4 a). Localized enrichment of Si along these decorated dislocation lines creates a rope-like appearance for dislocations with periodically varying widths along the length direction of those cylinders. Composition measurement by 1-D concentration profiles and proximity histograms yield comparable results. The Si concentration is ~ 14.4 ± 2.0 at.% at the dislocation core regions, corresponding to an enrichment of more than 20 times greater than the bulk Si concentration of 0.66 at.%. Meanwhile the average Si enrichment along the non-core regions of the dislocation lines is 5.7 ± 1.2 at.%. The average diameter of the Si-enriched zones around dislocation lines is 1.85 ± 0.4 nm. Crystallographic analysis on the grain constituting the APT needle provides greater insight into the Si decoration of dislocations, Fig. 4 b. The APT needle is aligned close to the {00 \(\:\stackrel{-}{2}\) } pole. All dislocations within the APT needle volume are located on { \(\:\stackrel{-}{1}2\) 1} planes, which are typical slip planes for the bcc system. Moreover, each dislocation has a 〈11 \(\:\stackrel{-}{1}\) 〉 slip direction parallel to the dislocation line segment (Fig. 4 b), indicating these dislocations are of screw character. Corroborating TEM analyses on the 1 dpa PM-HIP specimen reveals regions containing an array of piled-up dislocations, similar to APT observation, Fig. 4 c. A selected region (marked with a red dashed box) is shown at higher magnification in both bright field (BF) and dark field (DF). In the corresponding diffraction pattern of this selected region, the primary (brightly contrasting) diffraction spots represent the [012] zone axis of bcc Fe, while faint secondary diffraction spots are also present. DF TEM reveals an absence of precipitates or carbides in the selected area, indicating that the secondary diffraction spots must instead originate from low volume fraction phases that form along the array of piled-up dislocations. Indexing the secondary diffraction pattern reveals that these diffraction spots correspond to a Si-rich orthorhombic β-FeSi 2 phase, with \(\:{{\left[110\right]}_{\beta\:-FeSi}}_{2}\) as the zone axis, Fig. 4 c. These diffraction results, together with the high Si concentration at the dislocation cores, indicate that when Si radiation-induced segregation exceeds a critical concentration, dislocation core regions can transform into β-FeSi 2 phase linear complexions instead of remaining as Si-decorated dislocation lines. Further TEM analysis reveals that the \(\:\left\{1\stackrel{-}{2}1\right\}\) Fe plane trace coincides with the dislocation lines within the piled-up array, indicating the slip planes are \(\:\left\{1\stackrel{-}{2}1\right\}\) , consistent with APT observations. TEM projections of 〈111〉 bcc slip directions aligned along dislocation lines, further corroborate APT observations suggesting these dislocations are of screw character. Finding such a high density of screw dislocations in a bcc steel after 400ºC irradiation is surprising, as the mobility of screw dislocations is higher than that of edge dislocations. This is because the mobility of edge dislocation remains unchanged with temperature owing to the phonon effect, whilst the motion of screw dislocations is enhanced significantly through mechanisms like kink-pair formation at elevated temperatures 55 , 56 . Thus, the retention of screw dislocations in the 1 dpa PM-HIP specimen is likely due to the transformation of decorated dislocation cores into linear complexions of β-FeSi 2 . The resultant dislocation–complexion interactions create highly entangled networks and small angle grain boundaries, immobilizing screw dislocation segments in the 1 dpa PM-HIP specimen. Stepped dislocation segments, indicative of jog or kink-pair formations, are observed in PM-HIP 1 dpa, whereas these are sparse in cast 1 dpa and absent in both 3 dpa specimens, Fig. 5 . Discussion Formation of β-FeSi 2 linear complexions . Strong solute enrichment in combination with other factors can promote phase transformations. For example, extreme local elastic distortions can stimulate phase transformations of martensite to austenite through dislocation stress field-mediated austenite nucleation 21 , 29 . To elucidate formation mechanisms of β-FeSi 2 linear complexions along dislocation line segments, consider the phases in the Si-rich (i.e., 50–67% Si) region of the Fe-Si phase diagram 57 . There are two crystalline polymorphs of iron disilicide (FeSi 2 ): α-FeSi 2 high temperature phase stable over 960–1211ºC and having tetragonal lattice (space group P4/mmm), and β-FeSi 2 low temperature phase stable below 1002.5ºC with orthorhombic lattice (space group Cmca) 57 . Although the APT measured at dislocation core regions (~ 15 at.% Si) is inconsistent with stoichiometric β-FeSi 2 , experimentally obtained compositions of complexions often differ from the ideal stoichiometric concentration 24 . Precipitating Si-rich phases along dislocation cores may follow classical heterogeneous nucleation 58 , wherein the interaction between dislocation strain fields and Si solute atoms reduces the strain energy in presence of a precipitate embryo. Consequently, the solubility of Si in Fe decreases below equilibrium solubility, creating an opportunity for precipitation to initiate. Moreover, the screw character of dislocations (Fig. 4 b and c ) creates a favorable condition for Si segregation into dislocation cores. Unlike for edge dislocations, elastic distortions around screw dislocations contain no tensile or compressive components of hydrostatic pressure, creating a radially symmetric stress field 59 . Hence, screw dislocations have no specific preferential sites for solute segregation, leading to homogeneous Si segregation around screw dislocation cores (whereas edge dislocations will experience preferential solute segregation to the compressive side of the strain field 23 , 60 ). Even though β-FeSi 2 is present at dislocation core regions, these phases may first form as the high-temperature α-FeSi 2 phase within irradiation-induced thermal spikes. The formation of α-FeSi 2 in the PM-HIP G91 (i.e., bcc Fe lattice) is vacancy mediated. Unlike in conventional cast alloys, PM-HIP alloys contain retained porosity which reduces thermal conductivity and thus considerably alters phase stability and evolution 61 ; residual porosity is present in the investigated PM-HIP G91 ( Fig. S2 of supplementary file). Consequently, irradiation-induced thermal spikes can lead to localized incipient heating to temperatures in excess of the 960ºC necessary for the α-FeSi 2 transformation 57 . Irradiation-induced excess vacancies will then replace Fe and other substitutional solutes (e.g., Cr) on their lattice positions (notably, the concentration of Fe and major substitutional solute Cr are anti-correlated with Si concentration along dislocation lines, as shown in Fig. 4 a). As a result, the Si-enriched dislocations populated with excess vacancies could create a lattice as presented schematically in Fig. 6 . When the Fe sublattice contains ~ 12.5% excess vacancies 62 , the α-FeSi 2 structure can be visualized as a defected CsCl-type structure, wherein alternate layers of Fe along the c -direction (i.e., [001] Fe direction) are replaced by vacancies, Fig. 6 a. As a result, the axial ratio of c/2a is slightly less than unity, and the Si layers are slightly shifted towards each other, leading to the formation of α-FeSi 2 with lattice parameters a = 2.7047 Å, b = 2.7047 Å, and c = 5.1430 63 Å, Fig. 6 b. Subsequent formation of β-FeSi 2 is complex, requiring relocation of vacancies 64 . Further vacancy migration away from the body-centered position of the α-FeSi 2 lattice (Fig. 6 c ) , replacing Fe (like dislocation 1 and 2 in Fig. 4 a) or substitutional solute Cr (similar to dislocation 3 in Fig. 4 a) during irradiation can create an approximant β-FeSi 2 lattice, Fig. 6 d. Analogous to α-FeSi 2 , the β-FeSi 2 structure can be perceived as alternating layers of Fe atoms and vacancies along the [110] direction of the initial Fe lattice, Fig. 6 e. Then, to arrive at the final orthorhombic structure with lattice parameters a = 9.863 Å, b = 7.791 Å, and c = 7.833 Å, a series of complex atomic reshuffling occurs as described in refs. 65 . Implication of linear complexions in dislocation plasticity. The extended yielding phenomenon that occurs in the PM-HIP 1 dpa specimen is similar to static strain aging, which is typically observed in materials exhibiting Cottrell atmospheres that pin dislocations and require higher applied stresses for dislocations to break free from the solute environment. However, in the current PM-HIP 1 dpa specimens, the screw dislocations are unobstructed by substitutional solute or interstitial elements (typically C and N in steels) that typically comprise Cottrell atmospheres. Rather, screw dislocations are restrained by the transformed β-FeSi 2 phase within the dislocation core itself, which is a linear complexion 19 . The fluctuation in stress values and the discontinuous transition from the elastic into the elastic-plastic deformation regime during extended yielding is a result of successive pining and unpinning of dislocations at the linear complexions. Under an externally applied stress, dislocation bowing occurs in regions where the distance between linear complexions (i.e., pinning points, A and B in step 2 in Fig. 7 a) is largest 24 . The dislocation–complexion interaction energy will differ significantly from the dislocation–Si solute interaction energy in Fe, even though the chemical environments in both cases are similar (i.e., primarily entail Fe and Si interatomic potentials). The interaction between Si solute atoms and screw dislocations in a random solid solution Fe-9 at.% Si alloy requires a critical resolved shear stress (CRSS) of ~ 180 MPa at room temperature 66 . In contrast, for room temperature tensile deformation of PM-HIP 1 dpa steel, the interaction energy between screw dislocations and nanoscale β-FeSi 2 linear complexions will be quite different, due to crystallographic anisotropy-driven moduli. Variations in anisotropic factors ( A uvw ) and corresponding Hill bulk modulus ( B ) are calculated from elastic constants ( E uvw ) for different crystallographic orientations uvw are A 100 = 0.960, A 010 = 0.943, A 001 = 1.013, B a−axis = 464.8 GPa, B b−axis = 621.8 GPa, and B c−axis = 557.9 GPa 67 . This is fundamentally the primary difference between the traditional upper- and lower-yield point phenomenon associated with Cottrell atmospheres, compared to linear complexion mediated extended yielding. Additionally, the brittle nature of β-FeSi 2 ( K IC fracture toughness of 2–3 \(\:MPa\bullet\:\sqrt{{m}^{1/2}}\) , comparable to TiC 68 ) prevents other dislocations from shearing through the linear complexion. Prima facie, the dislocation bowing mechanism may seem similar to typical Orowan bowing during precipitate hardening, when a precipitate hinders the dislocation slip path and thus necessitating dislocations bow out to move past the obstacle. But bowing can only occur when shear stress is applied, enabling dislocations to overcome their line tension and take on a curvature. Unlike Orowan bowing, however, linear complexions provide no physical obstruction to dislocation motion, nor do they form new dislocation loops or segments 20 , 26 , 28 . Instead, linear complexions interact with dislocation stress fields and constrain their mobility. Further, while traditional precipitates lie on the same slip planes as dislocations, linear complexions are located on planes above or below the dislocation slip planes 24 , 25 . In the next phase of dislocation–linear complexion interactions, dislocation unpinning occurs gradually. Molecular dynamics (MD) and Monte Carlo (MC) simulations of linear complexions indicate that a dislocation from a dislocation dipole pair separates first, while the trailing dislocation remains attached to that particular nanoparticle (A and B in step 3 in Fig. 7 a) comprising the complexion 25 , 60 , 69 . As the bowing curvature increases, the dislocation from the dipole that remained attached now also separates from that nanoparticle, and the dipole becomes attached to the next nanoparticle (C and D in step 4 in Fig. 7 a) along the linear complexion. This sequence repeats as the unzipping process continues, as represented schematically in Fig. 7 a. Although this mechanism has not yet been directly validated experimentally, the present study reveals TEM evidence of dislocation dipole formation in PM-HIP 1 dpa G91 (Fig. 7 b); similar dislocation dipole formation due to interactions between straight screw dislocation segments and nanoparticles has been reported in Fe-Si systems 70 . At the very least, then, the presence of dislocation dipoles lends credence to the mechanism of dislocation–linear complexion nanoparticle interaction suggested in MD-MC simulations 25 , 69 . Two possible mechanisms have been proposed in the literature as operative for plasticity involving dislocation dipoles. The first is based on the formation of vacancy clusters and their diffusion through the lattice and along dislocation lines, contributing to dipole stabilization through climbing of its edge components 71 . The second mechanism initiates from dragging of jogs facilitated by the cross-slip of screw dislocations upon encountering obstacles during deformation 72 . Evidence herein of linear complexions along screw dislocations in the PM-HIP 1 dpa specimen, together with the aforementioned models which predict dislocation detachment through screw dislocation dipoles, suggest the jog-drag (second) mechanism plausibly governs initial deformation during extended yielding. Moreover, the high number density of β-FeSi 2 linear complexions provide favorable conditions for further dislocation pile-up during deformation, which can be attributed to the interaction between dislocation stress fields and linear complexions and ease of dislocation nucleation 73 . As a result, greater dislocation pile-up leads to higher localized stress, promoting cross-slip 74 . Moreover, this piled-up dislocation offers a conducive environment for generation of dislocation dipoles by interacting with existing screw dislocations 74 , and thus further supporting the jog-drag mechanism. Later stages of extended yielding, specifically as plastic strain accumulates without increases in stress, can be explained by nucleation of dislocations and evolution of the linear complexion structures. Because solute-decorated dislocations are stronger sinks for irradiation-induced defects, such defects will accumulate more readily near decorated dislocations, and thus continuously increase the stress built up around them. These solute decorated dislocations containing a large concentration of defects, remain immobile as they require significantly higher activation stress than regular dislocation motion 75 . Hence, the excess stress experienced by these decorated dislocations leads to rapid and significant dislocation multiplication and subsequent motion of newly formed dislocations 37 , resulting in a dislocation avalanche and strain burst as shown by discrete dislocation dynamics simulation 76 , thus enabling dislocation plasticity to proceed without more need of bulk applied stress. Another possibility is that the structure of linear complexion changes during deformation, specifically that β-FeSi 2 may transform to α-FeSi 2 . The β-FeSi 2 →α-FeSi 2 transformation can occur through either diffusion or deformation, the latter pathway being more energetically favorable 67 . The transformation energy for β-FeSi 2 →α-FeSi 2 transformation has a local minimum at 0.5 strain, when the β-FeSi 2 structure is under monoclinic (100)[001] strain as shown in Fig. 7 c. However, since α-FeSi 2 is not stable at room temperature 57 , the α-FeSi 2 phase within dislocation segments is metastable and will ultimately compromise the ability to form or maintain linear complexion configurations. Thus, the combination of dislocation multiplication from decorated dislocations and β-FeSi 2 →α-FeSi 2 structural changes inside linear complexions results in an end to linear complexion-mediated extended yielding. To conclude, this study has documented the irradiation-induced formation of linear complexions in Grade 91 ferritic steel, whereas linear complexions have previously only been documented in model alloy systems. The β-FeSi 2 complexion formation is made possible because the lower thermal conductivity of the powder-based (PM-HIP) Grade 91 steel enables retention of a high density of screw dislocations during alloy fabrication and irradiation. The observed β-FeSi 2 complexions provide unprecedented improvements in ductility after irradiation, through screw dislocation-mediated plasticity. The formation of these linear complexions enables the alloy to overcome the strength–ductility tradeoff, one of the most longstanding challenges that limit the service lifetime of nuclear structural alloys. These findings open new strategies for improving the mechanical performance of ferritic steels and other structural alloys commonly used in nuclear energy applications. More broadly, this work presents new directions for complexion engineering of future metallic alloys that are autonomously resilient to irradiation and offer unprecedented combinations of strength and ductility under irradiation. Methods Materials and irradiation . The materials selected for the current study are Grade 91 ferritic steels produced either by conventional casting or by powder metallurgy with hot isostatic pressing (PM-HIP) processes. The cast ingot and PM-HIP compact were provided by the Electric Power Research Institute (EPRI), USA. The PM-HIP was fabricated by consolidating gas atomized Grade 91 alloy powders under isostatic pressure of 103.4 MPa at a temperature of 1121°C for 4 hours. This was followed by a normalization heat treatment at 1060 ± 14°C for 2.5 hours, with subsequent forced air fan cooling. Afterward, it was tempered at 777 ± 14°C for 4.5 hours followed by air cooling. The cast ingot also underwent the same normalization and tempering treatments. The chemical compositions of the alloys were measured by inductively coupled plasma atomic emission spectroscopy (ICP-AES) and provided in Table S1 . Chemical analysis confirms the alloys conform to Grade 91 steel specifications in ASTM A387, and have compositions consistent with other Grade 91 steel process heats studied specifically for nuclear applications 46 . Both the PM-HIP and cast materials were further sectioned into circular discs and round tensile bars. Disc specimens had a diameter of 3 mm and a thickness of 0.15 mm, and were prepared using wire electrical discharge machining (EDM). Discs were then mechanically polished using SiC paper up to 1200 grit, followed by electropolishing in a solution of 10% perchloric acid and 90% methanol (by volume) at − 40°C and 35 V for 20 seconds. Round threaded tensile bars were prepared to ASTM E8 specifications having dimensions of 76.2 mm in length and 6.35 mm in gauge diameter. The tensile bars were machined at Idaho National Laboratory (INL) using a computer numerical control (CNC) machining to a surface roughness of 3.2 µm. Comprehensive details of specimen machining and preparation are provided in ref. 77 . For neutron irradiations, tensile bar and disc specimens were loaded into drop-in capsules, which act as a barrier between the primary reactor coolant and the specimens, and were located in the A-6, A-7, and A-8 positions in the Advanced Test Reactor (ATR) at INL. Comprehensive details of the capsule design, irradiation thermal and fluence calculations and measurements, and irradiation histories, are provided in ref. 77 . A total of four irradiated disc specimens and six irradiated tensile bars were selected for microstructural and mechanical characterization, with additional unirradiated specimens reserved as controls. The target irradiation temperature was 400ºC and target irradiation doses were 1 displacement per atom (dpa) and 3 dpa. Actual irradiation doses and temperatures for each specimen are listed in Table S2 . For the 1 dpa target specimens, the average dose was 1.00 ± 0.02 dpa at temperatures ranging 358–389°C, whilst for the 3 dpa target specimens, the average dose of 3.68 ± 0.23 dpa was achieved at temperatures 351–397°C. Mechanical testing and microstructural characterization. Quasi-static uniaxial tensile tests were conducted on irradiated tensile bar specimens, and on unirradiated reference specimens, following ASTM E8. The tests were performed at room temperature within an inert argon (Ar) atmosphere using a remote-operated Instron 5869 screw-driven load frame located inside the hot cells at the Hot Fuel Examination Facility (HFEF) Main Cell Window 13 M at the Materials and Fuels Complex (MFC), INL. Tensile tests were performed at an initial crosshead speed of 0.279 mm/min, corresponding to a strain rate of 8.78 × 10 − 3 s − 1 , until 10% engineering strain. Thereafter, crosshead speed was increased to 1.0 mm/min, i.e., strain rate of 3.15 × 10 − 2 s − 1 , until fracture. Microstructural characterization for the unirradiated and irradiated specimens was carried out on the disc specimens. Post-irradiation, the same electropolishing as was done prior to irradiation, was performed to remove minor surface damage resulting from handling, inspection, and decontamination. Focused ion beam (FIB) lift-out was conducted using a Thermo-Fisher Scientific (formerly FEI) Quanta 3D dual-beam scanning electron microscope (SEM), to extract lamellae for transmission electron microscopy (TEM) analyses, following established protocols 78 . To minimize FIB-induced damage, a protective Pt coating was deposited on the specimen surface prior to milling, mitigating Ga⁺ ion implantation and sputtering effects. A final polishing step using a 5 kV Ga⁺ ion beam, followed by 2 kV low-energy cleaning, was employed to mitigate surface damage and Ga⁺ implantation. TEM observations were performed using a Thermo-Fisher Scientific Tecnai TF30-FEG STwin TEM, operated at 300 kV, and equipped with a high-angle annular dark field (HAADF) detector for scanning transmission electron microscopy (STEM), located at the Microscopy and Characterization Suite (MaCS) at the Center for Advanced Energy Studies (CAES). Dislocations and dislocation loops were imaged using the down-zone bright field (BF) STEM technique. This technique facilitates dislocation loop imaging by relaxing diffraction conditions—specifically the g•b invisibility criterion—through the use of a convergent electron beam. This allows simultaneous visualization of loops with different Burgers vectors and orientations, significantly simplifying their quantification and enhancing imaging efficiency. Compared to traditional two-beam or weak-beam dark-field imaging, BF-STEM offers higher contrast, reduced acquisition time, and improved statistical accuracy 79 . Void characterization was conducted using through-focus BF TEM, in which voids appear as contrast-reversing features between under- and over-focused images. The average lamella thickness used for volumetric density calculations was determined using electron energy loss spectroscopy (EELS) by analyzing the zero-loss peak in energy-filtered TEM mode. Image datasets were processed using TEM Instrument Analysis (TIA) software and measurements were done with ImageJ software. The irradiated disc specimens were further used for atom probe tomography (APT) studies. APT needles were prepared using FIB lift-out and annular milling, on the same Thermo-Fisher Scientific Quanta 3D FIB/SEM at MaCS, CAES. Local electrode atom probe (LEAP) analyses were carried out using a Cameca LEAP 4000X HR at MaCS, CAES, operated in laser-pulsed mode at a base temperature of 45–50 K, 200 kHz pulse frequency, and 60 pJ laser energy. To ensure statistical confidence and reduce the influence of local heterogeneities in specimens, at least two needles per irradiation conditions were analyzed that exceeded 2 million ion detections per needle. APT data were reconstructed and analyzed using Cameca IVAS version 3.8.6 and AP Suite software. Declarations Acknowledgements The authors thank Dr. Donna Guillen, Jana Howard, and Jeremy Burgener of Idaho National Laboratory for their assistance with irradiation experiments and irradiated specimen handling; Dr. Zhongxia Shang of Purdue University for assistance with microscopy of unirradiated materials; Sri Sowmya Panuganti and Saquib Bin Habib of Purdue University for their assistance with data management; and Dr. Benjamin Sutton and David Gandy of the Electric Power Research Institute for material supply. 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Also discoverable on Platform About Our Team In Review Editorial Policies Advisory Board Help Center Resources Author Services Accessibility API Access RSS feed Manage Cookie Preferences © Research Square 2026 | ISSN 2693-5015 (online) Privacy Policy Terms of Service Do Not Sell My Personal Information {"props":{"pageProps":{"initialData":{"identity":"rs-7419483","acceptedTermsAndConditions":true,"allowDirectSubmit":false,"archivedVersions":[],"articleType":"Article","associatedPublications":[],"authors":[{"id":509782047,"identity":"0e3e96dd-d2c7-4dba-82ca-1789767f0ab3","order_by":0,"name":"Janelle Wharry","email":"data:image/png;base64,iVBORw0KGgoAAAANSUhEUgAAAZAAAAAyAQMAAABI0h/eAAAABlBMVEX///8AAABVwtN+AAAACXBIWXMAAA7EAAAOxAGVKw4bAAAAtUlEQVRIiWNgGAWjYBACxgYGxgcJBRY8YCYDAzNRWpgNEgwkSNACBGwSDAYSMA4RWphn95hVPDCQkDFnb27+wFBhndhA0GFzzpjdADnMsudgmwTDmXQitMzI3QbWYnAjsY2Bse0wcVoKoFqaPzD+I1ILA1RLgwRjAzFa5pz/LAHWcgbol4Rj6cYEtRjObkv8+KPCxt7gePvjDx9qrGUJa5mBzEsgpBwE5CUIqxkFo2AUjIKRDgBKgD04IVKfkAAAAABJRU5ErkJggg==","orcid":"","institution":"University of Illinois","correspondingAuthor":true,"prefix":"","firstName":"Janelle","middleName":"","lastName":"Wharry","suffix":""},{"id":509782048,"identity":"701276a8-e646-45d3-b12a-c3710b0f9794","order_by":1,"name":"Arya Chatterjee","email":"","orcid":"https://orcid.org/0000-0001-8250-6184","institution":"University of Illinois Urbana-Champaign","correspondingAuthor":false,"prefix":"","firstName":"Arya","middleName":"","lastName":"Chatterjee","suffix":""},{"id":509782049,"identity":"d75781b3-cfdc-4f93-9d0a-de9fe35c58cf","order_by":2,"name":"Soumita Mondal","email":"","orcid":"","institution":"University of Illinois Urbana-Champaign","correspondingAuthor":false,"prefix":"","firstName":"Soumita","middleName":"","lastName":"Mondal","suffix":""},{"id":509782050,"identity":"300dc168-6051-4363-a021-37bb0823af0c","order_by":3,"name":"Yu Lu","email":"","orcid":"","institution":"Eurofins","correspondingAuthor":false,"prefix":"","firstName":"Yu","middleName":"","lastName":"Lu","suffix":""},{"id":509782051,"identity":"44397d18-aa83-4425-b2c2-68e600f9845c","order_by":4,"name":"Yaqiao Wu","email":"","orcid":"https://orcid.org/0000-0002-5041-0935","institution":"Kansas State University","correspondingAuthor":false,"prefix":"","firstName":"Yaqiao","middleName":"","lastName":"Wu","suffix":""}],"badges":[],"createdAt":"2025-08-20 16:50:18","currentVersionCode":1,"declarations":"","doi":"10.21203/rs.3.rs-7419483/v1","doiUrl":"https://doi.org/10.21203/rs.3.rs-7419483/v1","draftVersion":[],"editorialEvents":[],"editorialNote":"","failedWorkflow":false,"files":[{"id":90565703,"identity":"927ef1a6-699e-46d6-9362-1a73316ad44a","added_by":"auto","created_at":"2025-09-04 07:20:50","extension":"png","order_by":1,"title":"Figure 1","display":"","copyAsset":false,"role":"figure","size":21670167,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cem\u003e(a) Engineering stress-strain curves of the investigated G91 steels from uniaxial tensile tests, with associated (b) fractographs of the unirradiated and irradiated cast and PM-HIP alloys, with example voids and quasi-cleavage regions indicated by red and yellow arrows, respectively. (c) Irradiated Grade 91 strength–ductility tradeoff (ductility quantified as % retained uniform elongation) from archival literature\u003c/em\u003e\u003csup\u003e39–42\u003c/sup\u003e\u003cem\u003e, showing that PM-HIP 1 dpa specimen from current study exhibits exceptional ductility.\u003c/em\u003e\u003c/p\u003e","description":"","filename":"Figure1.png","url":"https://assets-eu.researchsquare.com/files/rs-7419483/v1/c3fa10dbd93d71666fa95203.png"},{"id":90565432,"identity":"42164864-d59f-4c12-937d-fbbef5af97f0","added_by":"auto","created_at":"2025-09-04 07:12:50","extension":"png","order_by":2,"title":"Figure 2","display":"","copyAsset":false,"role":"figure","size":27289161,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cem\u003eTEM micrographs of (a-b) unirradiated cast and PM-HIP G91 exhibiting typical lath martensite structure with MX (black arrows) and M\u003c/em\u003e\u003csub\u003e\u003cem\u003e23\u003c/em\u003e\u003c/sub\u003e\u003cem\u003eC\u003c/em\u003e\u003csub\u003e\u003cem\u003e6\u003c/em\u003e\u003c/sub\u003e\u003cem\u003e precipitates (yellow arrows); (c-d) voids in 3 dpa irradiated cast and PM-HIP G91; (e-h) down-zone STEM micrographs of dislocation loops in cast and PM-HIP steels irradiated to (e-f) 1 dpa and (g-h) 3 dpa.\u003c/em\u003e\u003c/p\u003e","description":"","filename":"Figure2.png","url":"https://assets-eu.researchsquare.com/files/rs-7419483/v1/f90d9c4a7944e98241de63e8.png"},{"id":90565439,"identity":"05f22dcc-0076-4669-a44c-363d7a94cdeb","added_by":"auto","created_at":"2025-09-04 07:12:50","extension":"png","order_by":3,"title":"Figure 3","display":"","copyAsset":false,"role":"figure","size":79145918,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cem\u003e3D APT reconstructions of (a) 1 dpa cast, (b) 1 dpa PM-HIP, (c) 3 dpa cast and (d) 3 dpa PM-HIP Grade 91, showing irradiation-induced nanoclusters and segregation to dislocation lines. Isosurfaces are shown for Si concentrations of 3 at.%\u003c/em\u003e\u003c/p\u003e","description":"","filename":"Figure3.png","url":"https://assets-eu.researchsquare.com/files/rs-7419483/v1/3ddf017e71bef8c2e8619ea6.png"},{"id":90565427,"identity":"4755e972-cdb4-471a-aa8d-057383d7b5cd","added_by":"auto","created_at":"2025-09-04 07:12:49","extension":"png","order_by":4,"title":"Figure 4","display":"","copyAsset":false,"role":"figure","size":302987,"visible":true,"origin":"","legend":"\u003cp\u003eSee image above for figure legend\u003c/p\u003e","description":"","filename":"Figure4.png","url":"https://assets-eu.researchsquare.com/files/rs-7419483/v1/928971f5af8deef049dd7088.png"},{"id":90565428,"identity":"bfc3ebe5-c2db-44b5-8109-8d63a41a5782","added_by":"auto","created_at":"2025-09-04 07:12:50","extension":"png","order_by":5,"title":"Figure 5","display":"","copyAsset":false,"role":"figure","size":18673255,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cem\u003eTEM micrographs of revealing marginal and significant formation of stepped dislocation segments indicating jog and kink-pair formations in 1 dpa irradiated (a) cast and (b) PM-HIP samples, respectively. Whereas, 3 dpa irradiated (c) cast and (d) specimens do not indicate any such jog/kink-pair formations.\u003c/em\u003e\u003c/p\u003e","description":"","filename":"Figure5.png","url":"https://assets-eu.researchsquare.com/files/rs-7419483/v1/42196d15e68d0c8483bdd74f.png"},{"id":90564596,"identity":"c86f68f3-636c-463a-8ee9-29c6a24d5cb3","added_by":"auto","created_at":"2025-09-04 07:04:50","extension":"png","order_by":6,"title":"Figure 6","display":"","copyAsset":false,"role":"figure","size":15207063,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cem\u003eSchematic illustration of orthorhombic β-FeSi\u003c/em\u003e\u003csub\u003e\u003cem\u003e2\u003c/em\u003e\u003c/sub\u003e\u003cem\u003e transformation. (a) Fe bcc lattice with excess vacancies on alternating layers along [001]\u003c/em\u003e\u003csub\u003e\u003cem\u003eFe \u003c/em\u003e\u003c/sub\u003e\u003cem\u003edirection, can be visualized as (b) tetragonal structure of α-FeSi\u003c/em\u003e\u003csub\u003e\u003cem\u003e2\u003c/em\u003e\u003c/sub\u003e\u003cem\u003e. (c) Vacancy-mediated solute migration towards the body-centered position results in formation of (d) approximant β- FeSi\u003c/em\u003e\u003csub\u003e\u003cem\u003e2\u003c/em\u003e\u003c/sub\u003e\u003cem\u003e which can be visualized as (e) alternating layers of Fe atoms and vacancies along the [110]\u003c/em\u003e\u003csub\u003e\u003cem\u003eFe\u003c/em\u003e\u003c/sub\u003e\u003cem\u003e direction.\u003c/em\u003e\u003c/p\u003e","description":"","filename":"Figure6.png","url":"https://assets-eu.researchsquare.com/files/rs-7419483/v1/76a315007cee2e4f990fcf29.png"},{"id":90564622,"identity":"fcbb656e-ec16-46cc-9427-8235a92fc58a","added_by":"auto","created_at":"2025-09-04 07:04:51","extension":"png","order_by":7,"title":"Figure 7","display":"","copyAsset":false,"role":"figure","size":52649304,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cem\u003e(a) Schematic of bowing and unpinning process involving dislocation dipole and nanoparticles forming a linear complexion. (b) TEM micrographs indicating extensive dislocation dipole (red arrows) formation in 1 dpa irradiated PM-HIP specimen. (c) Schematic of atomic configuration showing deformation-induced phase transformation pathway for β-FeSi\u003c/em\u003e\u003csub\u003e\u003cem\u003e2\u003c/em\u003e\u003c/sub\u003e\u003cem\u003e→α-FeSi\u003c/em\u003e\u003csub\u003e\u003cem\u003e2\u003c/em\u003e\u003c/sub\u003e\u003cem\u003e transformation, where \u003c/em\u003e\u003cem\u003e\u003cstrong\u003ea\u003c/strong\u003e\u003c/em\u003e\u003cem\u003e, \u003c/em\u003e\u003cem\u003e\u003cstrong\u003eb\u003c/strong\u003e\u003c/em\u003e\u003cem\u003e and \u003c/em\u003e\u003cem\u003e\u003cstrong\u003ec\u003c/strong\u003e\u003c/em\u003e\u003cem\u003e refer to , \u0026nbsp;and \u0026nbsp;directions.\u003c/em\u003e\u003c/p\u003e","description":"","filename":"Figure7.png","url":"https://assets-eu.researchsquare.com/files/rs-7419483/v1/7aab6d2268682701d2e3d705.png"},{"id":90564581,"identity":"6ddbd1cb-4908-44f7-a570-d154512aa1b2","added_by":"auto","created_at":"2025-09-04 07:04:49","extension":"docx","order_by":1,"title":"","display":"","copyAsset":false,"role":"supplement","size":1011940,"visible":true,"origin":"","legend":"Supplementary file","description":"","filename":"Supplementaryfile.docx","url":"https://assets-eu.researchsquare.com/files/rs-7419483/v1/395533fbbb570b12a3774175.docx"}],"financialInterests":"There is \u003cb\u003eNO\u003c/b\u003e Competing Interest.","formattedTitle":"\u003cp\u003e\u003cstrong\u003eLinear complexions enable unprecedented ductility retention in neutron irradiated ferritic steel\u003c/strong\u003e\u003c/p\u003e","fulltext":[{"header":"Introduction","content":"\u003cp\u003eStrength and ductility are the key properties that influence the broad utilization of any structural material across varied applications and industries\u003csup\u003e\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e,\u003cspan citationid=\"CR2\" class=\"CitationRef\"\u003e2\u003c/span\u003e\u003c/sup\u003e, but these properties are often mutually exclusive \u003csup\u003e\u003cspan citationid=\"CR3\" class=\"CitationRef\"\u003e3\u003c/span\u003e\u003c/sup\u003e. The strength and ductility of metallic alloys, including those used for nuclear applications, depend on the presence of crystalline defects and their mobility under mechanical loading. The nature of plasticity in a metallic alloy is controlled by the interaction between dislocations (i.e., one-dimensional line defects) and all other types of defects including stacking faults\u003csup\u003e\u003cspan citationid=\"CR4\" class=\"CitationRef\"\u003e4\u003c/span\u003e\u003c/sup\u003e, grain boundaries\u003csup\u003e\u003cspan citationid=\"CR5\" class=\"CitationRef\"\u003e5\u003c/span\u003e\u003c/sup\u003e, precipitates\u003csup\u003e\u003cspan citationid=\"CR6\" class=\"CitationRef\"\u003e6\u003c/span\u003e\u003c/sup\u003e, solute atoms\u003csup\u003e\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e,\u003cspan citationid=\"CR8\" class=\"CitationRef\"\u003e8\u003c/span\u003e\u003c/sup\u003e, and other dislocations. Generally, as dislocations interact with defects, those defects can present strong obstacles that hinder the motion of dislocations, resulting in strengthening of the material. But this often comes at the sacrifice of the ability for the material to deform plastically, leading to a reduction in ductility.\u003c/p\u003e\u003cp\u003eThis strength\u0026ndash;ductility tradeoff becomes even more exacerbated in metallic alloys exposed to irradiation, such as during service in a nuclear reactor environment. Irradiation generates a supersaturation of point defects, which can evolve into extended defects such as clusters, dislocation loops, stacking faults, voids, solute segregation at the aforementioned features, precipitation, and twinning and phase transformations\u003csup\u003e\u003cspan citationid=\"CR9\" class=\"CitationRef\"\u003e9\u003c/span\u003e\u003c/sup\u003e. These irradiation-induced microstructures and microchemical gradients tend to hinder dislocation motion, resulting in an increase in strength and decrease in ductility. This phenomenon of \u003cem\u003eirradiation hardening and embrittlement\u003c/em\u003e poses a threat of brittle fracture in irradiated structural alloys, consequently compromising the safe operation of reactors.\u003c/p\u003e\u003cp\u003eRecently, numerous studies have sought to use defect engineering to intentionally introduce defects that can potentially help overcome the strength\u0026ndash;ductility tradeoff in metallic alloys. One such approach has been to extend plasticity by making twinning and mechanical phase transformations favorable via alloy design of twinning-induced plasticity (TWIP) and transformation-induced plasticity (TRIP) steels\u003csup\u003e\u003cspan citationid=\"CR2\" class=\"CitationRef\"\u003e2\u003c/span\u003e,\u003cspan citationid=\"CR10\" class=\"CitationRef\"\u003e10\u003c/span\u003e\u003c/sup\u003e. By obstructing dislocation motion, twins enhance strength, whilst the existence of multiple twins facilitates dislocation mobility through twin reorientation\u003csup\u003e\u003cspan citationid=\"CR2\" class=\"CitationRef\"\u003e2\u003c/span\u003e\u003c/sup\u003e. Similarly, phase transformations during deformation alter dislocation behavior in these new phases, often enhancing dislocation mobility\u003csup\u003e\u003cspan citationid=\"CR10\" class=\"CitationRef\"\u003e10\u003c/span\u003e\u003c/sup\u003e. In other cases, irradiation-induced defects such as voids and dislocation loops can also promote twinning and phase transformation in non-TWIP/TWIP alloys\u003csup\u003e\u003cspan additionalcitationids=\"CR12 CR13 CR14\" citationid=\"CR11\" class=\"CitationRef\"\u003e11\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR15\" class=\"CitationRef\"\u003e15\u003c/span\u003e\u003c/sup\u003e. Another approach to defect engineering is to take advantage of solute segregation, which is prevalent in irradiated materials. Chemical segregation is generally thought to hinder dislocation motion, thus contributing to hardening and loss of ductility\u003csup\u003e\u003cspan citationid=\"CR16\" class=\"CitationRef\"\u003e16\u003c/span\u003e,\u003cspan citationid=\"CR17\" class=\"CitationRef\"\u003e17\u003c/span\u003e\u003c/sup\u003e. However, if the segregation can be tailored such that it leads to the formation of a complexion, improvements in mechanical properties may be achieved\u003csup\u003e\u003cspan citationid=\"CR18\" class=\"CitationRef\"\u003e18\u003c/span\u003e\u003c/sup\u003e.\u003c/p\u003e\u003cp\u003eA linear complexion is a distinctly identifiable, but confined, chemical and structural state at dislocations\u003csup\u003e\u003cspan citationid=\"CR19\" class=\"CitationRef\"\u003e19\u003c/span\u003e\u003c/sup\u003e. Complexions can eventually give rise to stable\u003csup\u003e\u003cspan citationid=\"CR20\" class=\"CitationRef\"\u003e20\u003c/span\u003e\u003c/sup\u003e or metastable\u003csup\u003e\u003cspan citationid=\"CR21\" class=\"CitationRef\"\u003e21\u003c/span\u003e,\u003cspan citationid=\"CR22\" class=\"CitationRef\"\u003e22\u003c/span\u003e\u003c/sup\u003e phase formation. But if the extent of solute segregation to dislocation lines is insufficient to form stable precipitate growth and coalescence along those lines, linear arrays of nanoscale \u0026ldquo;precipitates\u0026rdquo; can form along the dislocation, creating what are known as linear complexions\u003csup\u003e\u003cspan additionalcitationids=\"CR20 CR21 CR22\" citationid=\"CR19\" class=\"CitationRef\"\u003e19\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR23\" class=\"CitationRef\"\u003e23\u003c/span\u003e\u003c/sup\u003e. These linear complexions can significantly influence strength and ductility and strength through dislocation pinning and unpinning\u003csup\u003e\u003cspan citationid=\"CR21\" class=\"CitationRef\"\u003e21\u003c/span\u003e,\u003cspan citationid=\"CR24\" class=\"CitationRef\"\u003e24\u003c/span\u003e,\u003cspan citationid=\"CR25\" class=\"CitationRef\"\u003e25\u003c/span\u003e\u003c/sup\u003e. The formation of linear complexions at dislocations has thus far primarily been predicted through atomistic simulations in model binary alloys such as Fe-Ni, Cu-Zr, Al-Cu, Ni-Al, and Al-Zr\u003csup\u003e\u003cspan citationid=\"CR20\" class=\"CitationRef\"\u003e20\u003c/span\u003e,\u003cspan additionalcitationids=\"CR23 CR24 CR25\" citationid=\"CR22\" class=\"CitationRef\"\u003e22\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR26\" class=\"CitationRef\"\u003e26\u003c/span\u003e\u003c/sup\u003e. Only a few studies have identified linear complexion formation experimentally by engineered thermo-mechanical treatments. For example, complexions can form in normalized Fe-9at.%Mn alloy through additional rolling to impart a high dislocation density, followed by extended tempering treatments allowing for Mn diffusion in Fe\u003csup\u003e\u003cspan citationid=\"CR19\" class=\"CitationRef\"\u003e19\u003c/span\u003e,\u003cspan citationid=\"CR21\" class=\"CitationRef\"\u003e21\u003c/span\u003e,\u003cspan citationid=\"CR27\" class=\"CitationRef\"\u003e27\u003c/span\u003e\u003c/sup\u003e. Nevertheless, nearly all reports of linear complexions are in model binary alloys\u003csup\u003e\u003cspan additionalcitationids=\"CR20 CR21 CR22 CR23 CR24 CR25 CR26 CR27 CR28 CR29\" citationid=\"CR19\" class=\"CitationRef\"\u003e19\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR30\" class=\"CitationRef\"\u003e30\u003c/span\u003e\u003c/sup\u003e.\u003c/p\u003e\u003cp\u003eHarnessing solute segregation to form linear complexions at dislocations, instead of forming clusters or precipitates, could thus remarkably improve ductility in an irradiation-hardened alloy. This pioneering study presents the novel formation of linear complexions at dislocations in an engineering alloy under service-relevant neutron irradiation conditions. Even more notably, these complexions promote excellent retention of ductility in irradiation-hardened materials, overcoming the longstanding strength\u0026ndash;ductility tradeoff. This work focuses on Grade 91 (G91) ferritic steel, used in current and advanced nuclear fission reactor structural components, manufactured using conventional casting and by powder metallurgy with hot isostatic pressing (PM-HIP). Materials are neutron irradiated to 1 and 3 displacements per atom (dpa) at 400\u0026ordm;C. The PM-HIP 1 dpa G91 forms linear complexions of the β-FeSi\u003csub\u003e2\u003c/sub\u003e phase along Si-segregated dislocation lines, enabling high retained ductility despite undergoing irradiation hardening. Coupling transmission electron microscopy (TEM) with atom probe tomography (APT) provides evidence of the structure and chemistry of the linear complexions. Finally, the manuscript discusses the plausible complexion formation and deformation mechanisms that enable commercial G91 ferritic steel to be amongst the first irradiated engineering alloys to overcome the strength-ductility tradeoff through complexion engineering.\u003c/p\u003e"},{"header":"Results","content":"\u003cp\u003e\u003cb\u003eMechanical behavior of unirradiated and irradiated G91 steel.\u003c/b\u003e Engineering stress-engineering strain curves from uniaxial tensile testing of the G91 steels are provided in Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ea and summarized in Table\u0026nbsp;\u003cspan refid=\"Tab1\" class=\"InternalRef\"\u003e1\u003c/span\u003e. Tensile tests reveal the general trend of irradiation hardening, as yield stress increases with neutron irradiation. Prior to irradiation, cast and PM-HIP G91 steels show similar yield strengths (YS) of 403 MPa and 447 MPa, ultimate tensile strengths (UTS) of 606 and 639 MPa, and uniform elongations (UE) of 10%, respectively. Upon 1 dpa neutron irradiation, PM-HIP exhibits higher hardening than cast (YS of 664 MPa compared to 628 MPa), similar UTS (737 MPa compared to 734 MPa), and retains higher ductility (UE of 9.1% as compared to 5.6%). At 3 dpa, the cast alloy exhibits higher hardening and strengthening than the PM-HIP (YS 765 MPa and UTS 819 MPa for PM-HIP; YS 801 MPa and UTS 867 MPa for cast), whilst both show a significant reduction in ductility, although PM-HIP continues to retain greater ductility than cast (UE of 4.4% and 2.3% for PM-HIP and cast, respectively). These superior irradiation hardening and embrittlement behaviors of the PM-HIP alloy as compared to its cast counterpart has also been observed in other ferritic steels and Ni-based alloys\u003csup\u003e\u003cspan citationid=\"CR31\" class=\"CitationRef\"\u003e31\u003c/span\u003e,\u003cspan citationid=\"CR32\" class=\"CitationRef\"\u003e32\u003c/span\u003e\u003c/sup\u003e Fractography of the broken tensile bars, Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003eb, reveals void-dominant dimpled fracture in unirradiated G91 steels, with some quasi-cleavage regions on the fracture surface of the unirradiated cast specimen. This quasi-cleavage fracture region becomes more prominent in the cast specimen after 1 dpa, whereas the PM-HIP retains dimpled fracture surfaces after 1 dpa, indicative of excellent retained ductility close to its pre-irradiated state. In both 3 dpa irradiated specimens, angular faceted features are present along with evidence of larger voids, consistent with the lower ductility.\u003c/p\u003e\u003cp\u003eThe mechanical properties of PM-HIP G91 steel after 1 dpa are noteworthy, as the alloy retains\u0026thinsp;\u0026gt;\u0026thinsp;90% of its pre-irradiation ductility, despite significant irradiation-induced strengthening, Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ec. A key driver for maintaining this exceptional ductility is the extended yielding phenomenon observed in 1 dpa irradiated PM-HIP G91, as depicted within the inset of Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ea, followed by gradual strain hardening. This extended yielding phenomenon is identified by the accumulation of ~\u0026thinsp;2% strain with overall negligible increase in bulk stress value, occurring between yielding and work hardening. These discontinuous yielding phenomena have been associated with various mechanisms: dislocation-solute interactions, such as Cottrell atmospheres\u003csup\u003e\u003cspan citationid=\"CR33\" class=\"CitationRef\"\u003e33\u003c/span\u003e\u003c/sup\u003e when upper and lower yield points are observed; L\u0026uuml;ders band formation in cases of strain localization; and dynamic strain aging effects, such as Portevin-Le Chatelier (PLC) behavior for the case of serrated or jerky flow\u003csup\u003e\u003cspan citationid=\"CR34\" class=\"CitationRef\"\u003e34\u003c/span\u003e\u003c/sup\u003e. Dislocation-solute interactions have primarily considered the effects of interstitial solute atoms (e.g., C and N) in Fe-based binary systems\u003csup\u003e\u003cspan citationid=\"CR33\" class=\"CitationRef\"\u003e33\u003c/span\u003e\u003c/sup\u003e, ferritic-martensitic alloys\u003csup\u003e\u003cspan citationid=\"CR35\" class=\"CitationRef\"\u003e35\u003c/span\u003e\u003c/sup\u003e, and austenitic Fe-Mn-C TWIP steels\u003csup\u003e\u003cspan citationid=\"CR36\" class=\"CitationRef\"\u003e36\u003c/span\u003e\u003c/sup\u003e. Similarly, under irradiation, when the rate of solute atom diffusion is slower than the dislocation glide velocity, L\u0026uuml;ders band phenomena can be observed through the formation of cleared channels in the microstructure\u003csup\u003e\u003cspan citationid=\"CR37\" class=\"CitationRef\"\u003e37\u003c/span\u003e\u003c/sup\u003e. By contrast, when solute diffusion rates are comparable to dislocation glide velocity, the PLC effect occurs through a repeated locking-unlocking mechanism between solutes and dislocations, and by propagating intermittent plasticity throughout the specimen\u003csup\u003e\u003cspan citationid=\"CR38\" class=\"CitationRef\"\u003e38\u003c/span\u003e\u003c/sup\u003e.\u003c/p\u003e\u003cp\u003e\u003cdiv class=\"gridtable\"\u003e\u003ctable float=\"Yes\" id=\"Tab1\" border=\"1\"\u003e\u003ccaption language=\"En\"\u003e\u003cdiv class=\"CaptionNumber\"\u003eTable 1\u003c/div\u003e\u003cdiv class=\"CaptionContent\"\u003e\u003cp\u003eMechanical properties of investigated irradiated G91 steels.\u003c/p\u003e\u003c/div\u003e\u003c/caption\u003e\u003ccolgroup cols=\"4\"\u003e\u003cdiv align=\"left\" class=\"colspec\" colname=\"c1\" colnum=\"1\"\u003e\u003c/div\u003e\u003cdiv align=\"char\" char=\".\" class=\"colspec\" colname=\"c2\" colnum=\"2\"\u003e\u003c/div\u003e\u003cdiv align=\"char\" char=\".\" class=\"colspec\" colname=\"c3\" colnum=\"3\"\u003e\u003c/div\u003e\u003cdiv align=\"char\" char=\".\" class=\"colspec\" colname=\"c4\" colnum=\"4\"\u003e\u003c/div\u003e\u003cthead\u003e\u003ctr\u003e\u003cth align=\"left\" colname=\"c1\"\u003e\u003cp\u003eMaterial\u003c/p\u003e\u003c/th\u003e\u003cth align=\"left\" colname=\"c2\"\u003e\u003cp\u003eYS (MPa)\u003c/p\u003e\u003c/th\u003e\u003cth align=\"left\" colname=\"c3\"\u003e\u003cp\u003eUTS (MPa)\u003c/p\u003e\u003c/th\u003e\u003cth align=\"left\" colname=\"c4\"\u003e\u003cp\u003eUniform elongation (%)\u003c/p\u003e\u003c/th\u003e\u003c/tr\u003e\u003c/thead\u003e\u003ctbody\u003e\u003ctr\u003e\u003ctd align=\"left\" colname=\"c1\"\u003e\u003cp\u003e\u003cb\u003eCast unirradiated\u003c/b\u003e\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c2\"\u003e\u003cp\u003e402.8\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e\u003cp\u003e606.0\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c4\"\u003e\u003cp\u003e9.88\u003c/p\u003e\u003c/td\u003e\u003c/tr\u003e\u003ctr\u003e\u003ctd align=\"left\" colname=\"c1\"\u003e\u003cp\u003e\u003cb\u003ePM-HIP unirradiated\u003c/b\u003e\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c2\"\u003e\u003cp\u003e447.2\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e\u003cp\u003e639.3\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c4\"\u003e\u003cp\u003e10.00\u003c/p\u003e\u003c/td\u003e\u003c/tr\u003e\u003ctr\u003e\u003ctd align=\"left\" colname=\"c1\"\u003e\u003cp\u003e\u003cb\u003eCast 1 dpa\u003c/b\u003e\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c2\"\u003e\u003cp\u003e627.8\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e\u003cp\u003e733.8\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c4\"\u003e\u003cp\u003e5.62\u003c/p\u003e\u003c/td\u003e\u003c/tr\u003e\u003ctr\u003e\u003ctd align=\"left\" colname=\"c1\"\u003e\u003cp\u003e\u003cb\u003ePM-HIP 1 dpa\u003c/b\u003e\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c2\"\u003e\u003cp\u003e663.7\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e\u003cp\u003e736.9\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c4\"\u003e\u003cp\u003e9.12\u003c/p\u003e\u003c/td\u003e\u003c/tr\u003e\u003ctr\u003e\u003ctd align=\"left\" colname=\"c1\"\u003e\u003cp\u003e\u003cb\u003eCast 3 dpa\u003c/b\u003e\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c2\"\u003e\u003cp\u003e800.9\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e\u003cp\u003e867.0\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c4\"\u003e\u003cp\u003e2.32\u003c/p\u003e\u003c/td\u003e\u003c/tr\u003e\u003ctr\u003e\u003ctd align=\"left\" colname=\"c1\"\u003e\u003cp\u003e\u003cb\u003ePM-HIP 3 dpa\u003c/b\u003e\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c2\"\u003e\u003cp\u003e764.8\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e\u003cp\u003e819.2\u003c/p\u003e\u003c/td\u003e\u003ctd align=\"char\" char=\".\" colname=\"c4\"\u003e\u003cp\u003e4.42\u003c/p\u003e\u003c/td\u003e\u003c/tr\u003e\u003c/tbody\u003e\u003c/colgroup\u003e\u003c/table\u003e\u003c/div\u003e\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003cp\u003e\u003cb\u003eCharacterization of microstructure to correlate mechanical properties\u003c/b\u003e. Transmission electron microscopic (TEM) analyses show that the initial microstructure of both cast and PM-HIP G91 steel is comprised of lath martensite with larger M\u003csub\u003e23\u003c/sub\u003eC\u003csub\u003e6\u003c/sub\u003e (M\u0026thinsp;=\u0026thinsp;Cr/Mo) type carbides and nanoscale microalloyed MX type precipitates (M\u0026thinsp;=\u0026thinsp;V, X\u0026thinsp;=\u0026thinsp;C/N), Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003e\u003cb\u003e(a, b\u003c/b\u003e). Such a microstructure is typical of 9Cr-1Mo F/M steels, consistently reported in numerous earlier studies\u003csup\u003e\u003cspan citationid=\"CR43\" class=\"CitationRef\"\u003e43\u003c/span\u003e,\u003cspan citationid=\"CR44\" class=\"CitationRef\"\u003e44\u003c/span\u003e\u003c/sup\u003e. Following irradiation, no cavities are present at ~\u0026thinsp;1 dpa, while only a negligible population of cavities is found at ~\u0026thinsp;3 dpa, Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003e\u003cb\u003e(c, d)\u003c/b\u003e. Cast G91 has ~\u0026thinsp;5 nm diameter cavities at a number density of 3.30\u0026times;10\u0026sup2;\u0026sup1; m\u003csup\u003e\u0026minus;\u0026thinsp;3\u003c/sup\u003e after 3 dpa, while the PM-HIP exhibits\u0026thinsp;~\u0026thinsp;4 nm diameter cavities at a number density of 2.13\u0026times;10\u0026sup2;\u0026sup1; m\u003csup\u003e\u0026minus;\u0026thinsp;3\u003c/sup\u003e. These cavities may result in only \u0026sim;0.001% volumetric swelling. The extent of cavity formation observed in this study is comparable to earlier work on Grade 91 steel from Tan, et al.\u003csup\u003e\u003cspan citationid=\"CR45\" class=\"CitationRef\"\u003e45\u003c/span\u003e\u003c/sup\u003e, wherein 4.36 dpa neutron irradiation at 469\u0026ordm;C results in an average cavity diameter of 3.0\u0026thinsp;\u0026plusmn;\u0026thinsp;1.1 nm (up to \u0026sim;6.8 nm) with a number density of (3.9\u0026thinsp;\u0026plusmn;\u0026thinsp;0.6) \u0026times; 10\u003csup\u003e21\u003c/sup\u003e m\u003csup\u003e\u0026minus;\u0026thinsp;3\u003c/sup\u003e, corresponding to ~\u0026thinsp;0.005% swelling. F/M steels are inherently resistant to void swelling, even at higher irradiation doses owing to their high density of martensite lath boundaries and extensive dislocation networks, as can be observed in Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003e\u003cb\u003e(a, b)\u003c/b\u003e. These lath boundaries and dislocation networks act as efficient sinks for recombination of point defects, thereby delaying the onset of void swelling\u003csup\u003e\u003cspan citationid=\"CR46\" class=\"CitationRef\"\u003e46\u003c/span\u003e\u003c/sup\u003e.\u003c/p\u003e\u003cp\u003eIrradiation-induced dislocation loops are populous in irradiated G91, Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003e\u003cb\u003e(e-h)\u003c/b\u003e. At 1 dpa, the average loop diameter is 27.8\u0026thinsp;\u0026plusmn;\u0026thinsp;5.1 nm for cast and 26.3\u0026thinsp;\u0026plusmn;\u0026thinsp;3.4 nm for PM-HIP, with corresponding number densities of 4.8 \u0026times; 10\u0026sup2;\u0026sup1; m\u003csup\u003e\u0026minus;\u0026thinsp;3\u003c/sup\u003e and 4.4 \u0026times; 10\u0026sup2;\u0026sup1; m\u003csup\u003e\u0026minus;\u0026thinsp;3\u003c/sup\u003e, respectively, \u003cb\u003eTable\u0026nbsp;2\u003c/b\u003e. Loops grow significantly by 3 dpa, reaching 43.5\u0026thinsp;\u0026plusmn;\u0026thinsp;3.7 nm and 38.3\u0026thinsp;\u0026plusmn;\u0026thinsp;2.8 nm, whilst their number densities decrease marginally to 4.07 \u0026times; 10\u0026sup2;\u0026sup1; m\u003csup\u003e\u0026minus;\u0026thinsp;3\u003c/sup\u003e and 3.7 \u0026times; 10\u0026sup2;\u0026sup1; m\u003csup\u003e\u0026minus;\u0026thinsp;3\u003c/sup\u003e for cast and PM-HIP steels, respectively, \u003cb\u003eTable\u0026nbsp;2\u003c/b\u003e. Loop sizes and densities are in agreement with previous reports on ~\u0026thinsp;350\u0026ndash;500\u0026ordm;C neutron irradiated F/M steels, in which loop diameters range 4.4\u0026ndash;6.1 nm at a density of 6.8\u0026ndash;7.8 \u0026times; 10\u003csup\u003e21\u003c/sup\u003e m\u003csup\u003e\u0026minus;\u0026thinsp;3\u003c/sup\u003e for doses 0.7\u0026ndash;1.1 dpa, and 10.6\u0026ndash;45.5 nm at a density of 0.8\u0026ndash;4.8 \u0026times; 10\u003csup\u003e21\u003c/sup\u003e m\u003csup\u003e\u0026minus;\u0026thinsp;3\u003c/sup\u003e for doses 3.9\u0026ndash;4.1 dpa\u003csup\u003e\u003cspan citationid=\"CR47\" class=\"CitationRef\"\u003e47\u003c/span\u003e,\u003cspan citationid=\"CR48\" class=\"CitationRef\"\u003e48\u003c/span\u003e\u003c/sup\u003e.\u003c/p\u003e\u003cp\u003eDislocation loops in 1 dpa cast and PM-HIP G91 are preferentially clustered near or along pre-existing dislocation lines and networks, Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003e\u003cb\u003e(e, f).\u003c/b\u003e By 3 dpa, loops appear more homogeneously distributed, though still exhibit some preferential location around martensite lath boundaries and subgrain boundaries, Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003e\u003cb\u003e(g, h)\u003c/b\u003e. Previous studies on low dose (generally\u0026thinsp;\u0026lt;\u0026thinsp;1 dpa) irradiated F/M steels also report similar heterogeneous distributions of loops near dislocation lines and networks\u003csup\u003e\u003cspan citationid=\"CR48\" class=\"CitationRef\"\u003e48\u003c/span\u003e\u003c/sup\u003e, indicating that dislocations may be sites for nucleation and trapping of loops. At higher dose levels, more vacancies and self-interstitial atoms (SIAs) are generated. Dislocation loops act as sinks for these point defects, absorbing them from the surrounding matrix\u003csup\u003e\u003cspan citationid=\"CR49\" class=\"CitationRef\"\u003e49\u003c/span\u003e\u003c/sup\u003e. Moreover, continuous absorption of SIAs leads to the growth of existing interstitial dislocation loops and can coalesce to other loops or dislocation network\u003csup\u003e\u003cspan citationid=\"CR50\" class=\"CitationRef\"\u003e50\u003c/span\u003e\u003c/sup\u003e, thereby increasing the loops size and becoming more uniformly distributed throughout the microstructure as observed at 3 dpa.\u003c/p\u003e\u003cp\u003eThe presence of loops along dislocation lines and networks in the 1 dpa specimens is similar to the irradiation defect\u0026ndash;dislocation interaction phenomenon often referred to as \u0026lsquo;rafts\u0026rsquo;\u003csup\u003e\u003cspan citationid=\"CR37\" class=\"CitationRef\"\u003e37\u003c/span\u003e\u003c/sup\u003e, wherein loops adhere to dislocation lines. Rafts form when irradiation-induced self-interstitial atom clusters migrate rapidly in a dislocation stress field, resulting in self-interstitial loops forming close to dislocations lines. These interstitial loops effectively pin dislocations and reduce their mobility. Only with additional stress, these locked dislocations can become mobile again. One may consider attributing the extended yielding phenomenon in 1 dpa PM-HIP G91 to dislocation\u0026ndash;raft locking\u0026ndash;unlocking, but this is unlikely. First, even though the loop\u0026ndash;dislocation distributions are similar in both cast and PM-HIP 1 dpa microstructures, the cast steel does not exhibit extended yielding. Second, rafts generally lead to strain localization by generating dislocation channels, which are absent in the irradiated-then-strained deformation microstructures collected near the fracture surface of all investigated specimens, \u003cb\u003eFig. \u003cspan refid=\"MOESM1\" class=\"InternalRef\"\u003eS1\u003c/span\u003e (supplementary file)\u003c/b\u003e.\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003cp\u003e\u003cb\u003eCharacterization of segregation\u003c/b\u003e. Irradiation-induced segregation of solute elements are characterized with atom probe tomography (APT). Reconstructed three dimensional (3D) atom maps obtained from the APT tips of the irradiated G91 specimens are depicted in Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003e\u003cb\u003e(a-d)\u003c/b\u003e. Discrete spherical solute clusters form in the cast specimen at 1 and 3 dpa, and in the PM-HIP specimen only at 3 dpa, Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003e\u003cb\u003e(a, c, d)\u003c/b\u003e; by contrast, in the PM-HIP specimen at 1 dpa, solute segregation occurs primarily along dislocation lines instead of nucleating discrete clusters, Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003eb. The primary solutes that form clusters are Ni, Mn, Si, and Cu, as listed in \u003cb\u003eTable\u0026nbsp;3\u003c/b\u003e. However, the dominant solute that decorates dislocation lines in the 1 dpa PM-HIP G91 is Si with sparse Ni clusters, Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003eb. Solute clustering is consistent with previous reports of irradiation-induced Ni-Mn-Si-rich (sometimes identified as stoichiometrically G-phase) and Cu-rich nanoclusters which form at number densities\u0026thinsp;~\u0026thinsp;10\u003csup\u003e23\u003c/sup\u003e m\u003csup\u003e\u0026minus;\u0026thinsp;3\u003c/sup\u003e over irradiation doses\u0026thinsp;~\u0026thinsp;1\u0026ndash;10 dpa at temperatures\u0026thinsp;~\u0026thinsp;300\u0026ndash;600\u0026ordm;C in F/M steels including Grade 91 type steels\u003csup\u003e\u003cspan citationid=\"CR47\" class=\"CitationRef\"\u003e47\u003c/span\u003e,\u003cspan citationid=\"CR48\" class=\"CitationRef\"\u003e48\u003c/span\u003e,\u003cspan citationid=\"CR51\" class=\"CitationRef\"\u003e51\u003c/span\u003e,\u003cspan citationid=\"CR52\" class=\"CitationRef\"\u003e52\u003c/span\u003e\u003c/sup\u003e. These clusters have previously been linked to irradiation hardening and are thus consistent with the observed uniaxial tensile stress-strain behavior of the cast 1 and 3 dpa and PM-HIP 3 dpa specimens herein. The reader is referred to previous studies\u003csup\u003e\u003cspan citationid=\"CR47\" class=\"CitationRef\"\u003e47\u003c/span\u003e,\u003cspan citationid=\"CR48\" class=\"CitationRef\"\u003e48\u003c/span\u003e,\u003cspan citationid=\"CR53\" class=\"CitationRef\"\u003e53\u003c/span\u003e,\u003cspan citationid=\"CR54\" class=\"CitationRef\"\u003e54\u003c/span\u003e\u003c/sup\u003e for details of nanocluster formation mechanisms and their implications on irradiation hardening. Here, this study will instead aim to explain the nature of solute segregation to dislocations in the 1 dpa PM-HIP specimen, and its implications on extended yielding.\u003c/p\u003e\u003cp\u003eTo characterize the irradiation-induced segregation to dislocations for the 1 dpa PM-HIP specimen, composition profiles are taken along and across the array of piled-up dislocation lines observed in the APT reconstruction Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003eb. Although Si enrichment is evident through the entirety of the dislocation pile-up array, the composition of Si fluctuates periodically along the dislocation lines; 1-D composition profiles reveal the period of Si fluctuation is ~\u0026thinsp;4.7\u0026thinsp;\u0026plusmn;\u0026thinsp;1.2 nm, Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ea. The Si concentration appears anti-correlated with either the Fe or Cr concentration, suggesting a substitutional solute exchange formation mechanism (profiles 1\u0026ndash;3, Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ea). Localized enrichment of Si along these decorated dislocation lines creates a rope-like appearance for dislocations with periodically varying widths along the length direction of those cylinders. Composition measurement by 1-D concentration profiles and proximity histograms yield comparable results. The Si concentration is ~\u0026thinsp;14.4\u0026thinsp;\u0026plusmn;\u0026thinsp;2.0 at.% at the dislocation core regions, corresponding to an enrichment of more than 20 times greater than the bulk Si concentration of 0.66 at.%. Meanwhile the average Si enrichment along the non-core regions of the dislocation lines is 5.7\u0026thinsp;\u0026plusmn;\u0026thinsp;1.2 at.%. The average diameter of the Si-enriched zones around dislocation lines is 1.85\u0026thinsp;\u0026plusmn;\u0026thinsp;0.4 nm. Crystallographic analysis on the grain constituting the APT needle provides greater insight into the Si decoration of dislocations, Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003eb. The APT needle is aligned close to the {00\u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:\\stackrel{-}{2}\\)\u003c/span\u003e\u003c/span\u003e} pole. All dislocations within the APT needle volume are located on {\u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:\\stackrel{-}{1}2\\)\u003c/span\u003e\u003c/span\u003e1} planes, which are typical slip planes for the bcc system. Moreover, each dislocation has a 〈11\u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:\\stackrel{-}{1}\\)\u003c/span\u003e\u003c/span\u003e〉 slip direction parallel to the dislocation line segment (Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003eb), indicating these dislocations are of screw character.\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003cp\u003eCorroborating TEM analyses on the 1 dpa PM-HIP specimen reveals regions containing an array of piled-up dislocations, similar to APT observation, Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ec. A selected region (marked with a red dashed box) is shown at higher magnification in both bright field (BF) and dark field (DF). In the corresponding diffraction pattern of this selected region, the primary (brightly contrasting) diffraction spots represent the [012] zone axis of bcc Fe, while faint secondary diffraction spots are also present. DF TEM reveals an absence of precipitates or carbides in the selected area, indicating that the secondary diffraction spots must instead originate from low volume fraction phases that form along the array of piled-up dislocations. Indexing the secondary diffraction pattern reveals that these diffraction spots correspond to a Si-rich orthorhombic β-FeSi\u003csub\u003e2\u003c/sub\u003e phase, with \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:{{\\left[110\\right]}_{\\beta\\:-FeSi}}_{2}\\)\u003c/span\u003e\u003c/span\u003e as the zone axis, Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ec. These diffraction results, together with the high Si concentration at the dislocation cores, indicate that when Si radiation-induced segregation exceeds a critical concentration, dislocation core regions can transform into β-FeSi\u003csub\u003e2\u003c/sub\u003e phase linear complexions instead of remaining as Si-decorated dislocation lines.\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003cp\u003eFurther TEM analysis reveals that the \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:\\left\\{1\\stackrel{-}{2}1\\right\\}\\)\u003c/span\u003e\u003c/span\u003e\u003csub\u003eFe\u003c/sub\u003e plane trace coincides with the dislocation lines within the piled-up array, indicating the slip planes are \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:\\left\\{1\\stackrel{-}{2}1\\right\\}\\)\u003c/span\u003e\u003c/span\u003e, consistent with APT observations. TEM projections of 〈111〉\u003csub\u003ebcc\u003c/sub\u003e slip directions aligned along dislocation lines, further corroborate APT observations suggesting these dislocations are of screw character. Finding such a high density of screw dislocations in a bcc steel after 400\u0026ordm;C irradiation is surprising, as the mobility of screw dislocations is higher than that of edge dislocations. This is because the mobility of edge dislocation remains unchanged with temperature owing to the phonon effect, whilst the motion of screw dislocations is enhanced significantly through mechanisms like kink-pair formation at elevated temperatures\u003csup\u003e\u003cspan citationid=\"CR55\" class=\"CitationRef\"\u003e55\u003c/span\u003e,\u003cspan citationid=\"CR56\" class=\"CitationRef\"\u003e56\u003c/span\u003e\u003c/sup\u003e. Thus, the retention of screw dislocations in the 1 dpa PM-HIP specimen is likely due to the transformation of decorated dislocation cores into linear complexions of β-FeSi\u003csub\u003e2\u003c/sub\u003e. The resultant dislocation\u0026ndash;complexion interactions create highly entangled networks and small angle grain boundaries, immobilizing screw dislocation segments in the 1 dpa PM-HIP specimen. Stepped dislocation segments, indicative of jog or kink-pair formations, are observed in PM-HIP 1 dpa, whereas these are sparse in cast 1 dpa and absent in both 3 dpa specimens, Fig.\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003e.\u003c/p\u003e\u003cp\u003e\u003c/p\u003e"},{"header":"Discussion","content":"\u003cp\u003e\u003cb\u003eFormation of β-FeSi\u003c/b\u003e\u003csub\u003e\u003cb\u003e2\u003c/b\u003e\u003c/sub\u003e \u003cb\u003elinear complexions\u003c/b\u003e. Strong solute enrichment in combination with other factors can promote phase transformations. For example, extreme local elastic distortions can stimulate phase transformations of martensite to austenite through dislocation stress field-mediated austenite nucleation\u003csup\u003e\u003cspan citationid=\"CR21\" class=\"CitationRef\"\u003e21\u003c/span\u003e,\u003cspan citationid=\"CR29\" class=\"CitationRef\"\u003e29\u003c/span\u003e\u003c/sup\u003e. To elucidate formation mechanisms of β-FeSi\u003csub\u003e2\u003c/sub\u003e linear complexions along dislocation line segments, consider the phases in the Si-rich (i.e., 50\u0026ndash;67% Si) region of the Fe-Si phase diagram\u003csup\u003e\u003cspan citationid=\"CR57\" class=\"CitationRef\"\u003e57\u003c/span\u003e\u003c/sup\u003e. There are two crystalline polymorphs of iron disilicide (FeSi\u003csub\u003e2\u003c/sub\u003e): α-FeSi\u003csub\u003e2\u003c/sub\u003e high temperature phase stable over 960\u0026ndash;1211\u0026ordm;C and having tetragonal lattice (space group P4/mmm), and β-FeSi\u003csub\u003e2\u003c/sub\u003e low temperature phase stable below 1002.5\u0026ordm;C with orthorhombic lattice (space group Cmca)\u003csup\u003e\u003cspan citationid=\"CR57\" class=\"CitationRef\"\u003e57\u003c/span\u003e\u003c/sup\u003e. Although the APT measured at dislocation core regions (~\u0026thinsp;15 at.% Si) is inconsistent with stoichiometric β-FeSi\u003csub\u003e2\u003c/sub\u003e, experimentally obtained compositions of complexions often differ from the ideal stoichiometric concentration\u003csup\u003e\u003cspan citationid=\"CR24\" class=\"CitationRef\"\u003e24\u003c/span\u003e\u003c/sup\u003e.\u003c/p\u003e\u003cp\u003ePrecipitating Si-rich phases along dislocation cores may follow classical heterogeneous nucleation\u003csup\u003e\u003cspan citationid=\"CR58\" class=\"CitationRef\"\u003e58\u003c/span\u003e\u003c/sup\u003e, wherein the interaction between dislocation strain fields and Si solute atoms reduces the strain energy in presence of a precipitate embryo. Consequently, the solubility of Si in Fe decreases below equilibrium solubility, creating an opportunity for precipitation to initiate. Moreover, the screw character of dislocations (Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003eb and \u003cb\u003ec\u003c/b\u003e) creates a favorable condition for Si segregation into dislocation cores. Unlike for edge dislocations, elastic distortions around screw dislocations contain no tensile or compressive components of hydrostatic pressure, creating a radially symmetric stress field\u003csup\u003e\u003cspan citationid=\"CR59\" class=\"CitationRef\"\u003e59\u003c/span\u003e\u003c/sup\u003e. Hence, screw dislocations have no specific preferential sites for solute segregation, leading to homogeneous Si segregation around screw dislocation cores (whereas edge dislocations will experience preferential solute segregation to the compressive side of the strain field\u003csup\u003e\u003cspan citationid=\"CR23\" class=\"CitationRef\"\u003e23\u003c/span\u003e,\u003cspan citationid=\"CR60\" class=\"CitationRef\"\u003e60\u003c/span\u003e\u003c/sup\u003e).\u003c/p\u003e\u003cp\u003eEven though β-FeSi\u003csub\u003e2\u003c/sub\u003e is present at dislocation core regions, these phases may first form as the high-temperature α-FeSi\u003csub\u003e2\u003c/sub\u003e phase within irradiation-induced thermal spikes. The formation of α-FeSi\u003csub\u003e2\u003c/sub\u003e in the PM-HIP G91 (i.e., bcc Fe lattice) is vacancy mediated. Unlike in conventional cast alloys, PM-HIP alloys contain retained porosity which reduces thermal conductivity and thus considerably alters phase stability and evolution\u003csup\u003e\u003cspan citationid=\"CR61\" class=\"CitationRef\"\u003e61\u003c/span\u003e\u003c/sup\u003e; residual porosity is present in the investigated PM-HIP G91 (\u003cb\u003eFig. S2\u003c/b\u003e of supplementary file). Consequently, irradiation-induced thermal spikes can lead to localized incipient heating to temperatures in excess of the 960\u0026ordm;C necessary for the α-FeSi\u003csub\u003e2\u003c/sub\u003e transformation\u003csup\u003e\u003cspan citationid=\"CR57\" class=\"CitationRef\"\u003e57\u003c/span\u003e\u003c/sup\u003e. Irradiation-induced excess vacancies will then replace Fe and other substitutional solutes (e.g., Cr) on their lattice positions (notably, the concentration of Fe and major substitutional solute Cr are anti-correlated with Si concentration along dislocation lines, as shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ea). As a result, the Si-enriched dislocations populated with excess vacancies could create a lattice as presented schematically in Fig.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003e. When the Fe sublattice contains\u0026thinsp;~\u0026thinsp;12.5% excess vacancies\u003csup\u003e\u003cspan citationid=\"CR62\" class=\"CitationRef\"\u003e62\u003c/span\u003e\u003c/sup\u003e, the α-FeSi\u003csub\u003e2\u003c/sub\u003e structure can be visualized as a defected CsCl-type structure, wherein alternate layers of Fe along the \u003cem\u003ec\u003c/em\u003e-direction (i.e., [001]\u003csub\u003eFe\u003c/sub\u003e direction) are replaced by vacancies, Fig.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003ea. As a result, the axial ratio of \u003cem\u003ec/2a\u003c/em\u003e is slightly less than unity, and the Si layers are slightly shifted towards each other, leading to the formation of α-FeSi\u003csub\u003e2\u003c/sub\u003e with lattice parameters \u003cem\u003ea\u003c/em\u003e\u0026thinsp;=\u0026thinsp;2.7047 \u0026Aring;, \u003cem\u003eb\u003c/em\u003e\u0026thinsp;=\u0026thinsp;2.7047 \u0026Aring;, and \u003cem\u003ec\u003c/em\u003e\u0026thinsp;=\u0026thinsp;5.1430\u003csup\u003e63\u003c/sup\u003e \u0026Aring;, Fig.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003eb. Subsequent formation of β-FeSi\u003csub\u003e2\u003c/sub\u003e is complex, requiring relocation of vacancies\u003csup\u003e\u003cspan citationid=\"CR64\" class=\"CitationRef\"\u003e64\u003c/span\u003e\u003c/sup\u003e. Further vacancy migration away from the body-centered position of the α-FeSi\u003csub\u003e2\u003c/sub\u003e lattice (Fig.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003ec\u003cb\u003e)\u003c/b\u003e, replacing Fe (like dislocation \u003cb\u003e1\u003c/b\u003e and \u003cb\u003e2\u003c/b\u003e in Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ea) or substitutional solute Cr (similar to dislocation \u003cb\u003e3\u003c/b\u003e in Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ea) during irradiation can create an approximant β-FeSi\u003csub\u003e2\u003c/sub\u003e lattice, Fig.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003ed. Analogous to α-FeSi\u003csub\u003e2\u003c/sub\u003e, the β-FeSi\u003csub\u003e2\u003c/sub\u003e structure can be perceived as alternating layers of Fe atoms and vacancies along the [110] direction of the initial Fe lattice, Fig.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003ee. Then, to arrive at the final orthorhombic structure with lattice parameters \u003cem\u003ea\u003c/em\u003e\u0026thinsp;=\u0026thinsp;9.863 \u0026Aring;, \u003cem\u003eb\u003c/em\u003e\u0026thinsp;=\u0026thinsp;7.791 \u0026Aring;, and \u003cem\u003ec\u003c/em\u003e\u0026thinsp;=\u0026thinsp;7.833 \u0026Aring;, a series of complex atomic reshuffling occurs as described in refs.\u003csup\u003e\u003cspan citationid=\"CR65\" class=\"CitationRef\"\u003e65\u003c/span\u003e\u003c/sup\u003e.\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003cp\u003e\u003cb\u003eImplication of linear complexions in dislocation plasticity.\u003c/b\u003e The extended yielding phenomenon that occurs in the PM-HIP 1 dpa specimen is similar to static strain aging, which is typically observed in materials exhibiting Cottrell atmospheres that pin dislocations and require higher applied stresses for dislocations to break free from the solute environment. However, in the current PM-HIP 1 dpa specimens, the screw dislocations are unobstructed by substitutional solute or interstitial elements (typically C and N in steels) that typically comprise Cottrell atmospheres. Rather, screw dislocations are restrained by the transformed β-FeSi\u003csub\u003e2\u003c/sub\u003e phase within the dislocation core itself, which is a linear complexion\u003csup\u003e\u003cspan citationid=\"CR19\" class=\"CitationRef\"\u003e19\u003c/span\u003e\u003c/sup\u003e.\u003c/p\u003e\u003cp\u003eThe fluctuation in stress values and the discontinuous transition from the elastic into the elastic-plastic deformation regime during extended yielding is a result of successive pining and unpinning of dislocations at the linear complexions. Under an externally applied stress, dislocation bowing occurs in regions where the distance between linear complexions (i.e., pinning points, A and B in step 2 in Fig.\u0026nbsp;\u003cspan refid=\"Fig7\" class=\"InternalRef\"\u003e7\u003c/span\u003ea) is largest\u003csup\u003e\u003cspan citationid=\"CR24\" class=\"CitationRef\"\u003e24\u003c/span\u003e\u003c/sup\u003e. The dislocation\u0026ndash;complexion interaction energy will differ significantly from the dislocation\u0026ndash;Si solute interaction energy in Fe, even though the chemical environments in both cases are similar (i.e., primarily entail Fe and Si interatomic potentials). The interaction between Si solute atoms and screw dislocations in a random solid solution Fe-9 at.% Si alloy requires a critical resolved shear stress (CRSS) of ~\u0026thinsp;180 MPa at room temperature\u003csup\u003e\u003cspan citationid=\"CR66\" class=\"CitationRef\"\u003e66\u003c/span\u003e\u003c/sup\u003e. In contrast, for room temperature tensile deformation of PM-HIP 1 dpa steel, the interaction energy between screw dislocations and nanoscale β-FeSi\u003csub\u003e2\u003c/sub\u003e linear complexions will be quite different, due to crystallographic anisotropy-driven moduli. Variations in anisotropic factors (\u003cem\u003eA\u003c/em\u003e\u003csub\u003e\u003cem\u003euvw\u003c/em\u003e\u003c/sub\u003e) and corresponding Hill bulk modulus (\u003cem\u003eB\u003c/em\u003e) are calculated from elastic constants (\u003cem\u003eE\u003c/em\u003e\u003csub\u003e\u003cem\u003euvw\u003c/em\u003e\u003c/sub\u003e) for different crystallographic orientations \u003cem\u003euvw\u003c/em\u003e are \u003cem\u003eA\u003c/em\u003e\u003csub\u003e\u003cem\u003e100\u003c/em\u003e\u003c/sub\u003e\u0026thinsp;=\u0026thinsp;0.960, \u003cem\u003eA\u003c/em\u003e\u003csub\u003e\u003cem\u003e010\u003c/em\u003e\u003c/sub\u003e\u0026thinsp;=\u0026thinsp;0.943, \u003cem\u003eA\u003c/em\u003e\u003csub\u003e\u003cem\u003e001\u003c/em\u003e\u003c/sub\u003e\u0026thinsp;=\u0026thinsp;1.013, \u003cem\u003eB\u003c/em\u003e\u003csub\u003e\u003cem\u003ea\u0026minus;axis\u003c/em\u003e\u003c/sub\u003e = 464.8 GPa, \u003cem\u003eB\u003c/em\u003e\u003csub\u003e\u003cem\u003eb\u0026minus;axis\u003c/em\u003e\u003c/sub\u003e = 621.8 GPa, and \u003cem\u003eB\u003c/em\u003e\u003csub\u003e\u003cem\u003ec\u0026minus;axis\u003c/em\u003e\u003c/sub\u003e = 557.9 GPa\u003csup\u003e\u003cspan citationid=\"CR67\" class=\"CitationRef\"\u003e67\u003c/span\u003e\u003c/sup\u003e. This is fundamentally the primary difference between the traditional upper- and lower-yield point phenomenon associated with Cottrell atmospheres, compared to linear complexion mediated extended yielding. Additionally, the brittle nature of β-FeSi\u003csub\u003e2\u003c/sub\u003e (\u003cem\u003eK\u003c/em\u003e\u003csub\u003e\u003cem\u003eIC\u003c/em\u003e\u003c/sub\u003e fracture toughness of 2\u0026ndash;3 \u003cspan class=\"InlineEquation\"\u003e\u003cspan class=\"mathinline\"\u003e\\(\\:MPa\\bullet\\:\\sqrt{{m}^{1/2}}\\)\u003c/span\u003e\u003c/span\u003e, comparable to TiC\u003csup\u003e\u003cspan citationid=\"CR68\" class=\"CitationRef\"\u003e68\u003c/span\u003e\u003c/sup\u003e) prevents other dislocations from shearing through the linear complexion. Prima facie, the dislocation bowing mechanism may seem similar to typical Orowan bowing during precipitate hardening, when a precipitate hinders the dislocation slip path and thus necessitating dislocations bow out to move past the obstacle. But bowing can only occur when shear stress is applied, enabling dislocations to overcome their line tension and take on a curvature. Unlike Orowan bowing, however, linear complexions provide no physical obstruction to dislocation motion, nor do they form new dislocation loops or segments\u003csup\u003e\u003cspan citationid=\"CR20\" class=\"CitationRef\"\u003e20\u003c/span\u003e,\u003cspan citationid=\"CR26\" class=\"CitationRef\"\u003e26\u003c/span\u003e,\u003cspan citationid=\"CR28\" class=\"CitationRef\"\u003e28\u003c/span\u003e\u003c/sup\u003e. Instead, linear complexions interact with dislocation stress fields and constrain their mobility. Further, while traditional precipitates lie on the same slip planes as dislocations, linear complexions are located on planes above or below the dislocation slip planes\u003csup\u003e\u003cspan citationid=\"CR24\" class=\"CitationRef\"\u003e24\u003c/span\u003e,\u003cspan citationid=\"CR25\" class=\"CitationRef\"\u003e25\u003c/span\u003e\u003c/sup\u003e.\u003c/p\u003e\u003cp\u003eIn the next phase of dislocation\u0026ndash;linear complexion interactions, dislocation unpinning occurs gradually. Molecular dynamics (MD) and Monte Carlo (MC) simulations of linear complexions indicate that a dislocation from a dislocation dipole pair separates first, while the trailing dislocation remains attached to that particular nanoparticle (A and B in step 3 in Fig.\u0026nbsp;\u003cspan refid=\"Fig7\" class=\"InternalRef\"\u003e7\u003c/span\u003ea) comprising the complexion\u003csup\u003e\u003cspan citationid=\"CR25\" class=\"CitationRef\"\u003e25\u003c/span\u003e,\u003cspan citationid=\"CR60\" class=\"CitationRef\"\u003e60\u003c/span\u003e,\u003cspan citationid=\"CR69\" class=\"CitationRef\"\u003e69\u003c/span\u003e\u003c/sup\u003e. As the bowing curvature increases, the dislocation from the dipole that remained attached now also separates from that nanoparticle, and the dipole becomes attached to the next nanoparticle (C and D in step 4 in Fig.\u0026nbsp;\u003cspan refid=\"Fig7\" class=\"InternalRef\"\u003e7\u003c/span\u003ea) along the linear complexion. This sequence repeats as the unzipping process continues, as represented schematically in Fig.\u0026nbsp;\u003cspan refid=\"Fig7\" class=\"InternalRef\"\u003e7\u003c/span\u003ea. Although this mechanism has not yet been directly validated experimentally, the present study reveals TEM evidence of dislocation dipole formation in PM-HIP 1 dpa G91 (Fig.\u0026nbsp;\u003cspan refid=\"Fig7\" class=\"InternalRef\"\u003e7\u003c/span\u003eb); similar dislocation dipole formation due to interactions between straight screw dislocation segments and nanoparticles has been reported in Fe-Si systems\u003csup\u003e\u003cspan citationid=\"CR70\" class=\"CitationRef\"\u003e70\u003c/span\u003e\u003c/sup\u003e. At the very least, then, the presence of dislocation dipoles lends credence to the mechanism of dislocation\u0026ndash;linear complexion nanoparticle interaction suggested in MD-MC simulations\u003csup\u003e\u003cspan citationid=\"CR25\" class=\"CitationRef\"\u003e25\u003c/span\u003e,\u003cspan citationid=\"CR69\" class=\"CitationRef\"\u003e69\u003c/span\u003e\u003c/sup\u003e.\u003c/p\u003e\u003cp\u003eTwo possible mechanisms have been proposed in the literature as operative for plasticity involving dislocation dipoles. The first is based on the formation of vacancy clusters and their diffusion through the lattice and along dislocation lines, contributing to dipole stabilization through climbing of its edge components\u003csup\u003e\u003cspan citationid=\"CR71\" class=\"CitationRef\"\u003e71\u003c/span\u003e\u003c/sup\u003e. The second mechanism initiates from dragging of jogs facilitated by the cross-slip of screw dislocations upon encountering obstacles during deformation\u003csup\u003e\u003cspan citationid=\"CR72\" class=\"CitationRef\"\u003e72\u003c/span\u003e\u003c/sup\u003e. Evidence herein of linear complexions along screw dislocations in the PM-HIP 1 dpa specimen, together with the aforementioned models which predict dislocation detachment through screw dislocation dipoles, suggest the jog-drag (second) mechanism plausibly governs initial deformation during extended yielding. Moreover, the high number density of β-FeSi\u003csub\u003e2\u003c/sub\u003e linear complexions provide favorable conditions for further dislocation pile-up during deformation, which can be attributed to the interaction between dislocation stress fields and linear complexions and ease of dislocation nucleation\u003csup\u003e\u003cspan citationid=\"CR73\" class=\"CitationRef\"\u003e73\u003c/span\u003e\u003c/sup\u003e. As a result, greater dislocation pile-up leads to higher localized stress, promoting cross-slip\u003csup\u003e\u003cspan citationid=\"CR74\" class=\"CitationRef\"\u003e74\u003c/span\u003e\u003c/sup\u003e. Moreover, this piled-up dislocation offers a conducive environment for generation of dislocation dipoles by interacting with existing screw dislocations\u003csup\u003e\u003cspan citationid=\"CR74\" class=\"CitationRef\"\u003e74\u003c/span\u003e\u003c/sup\u003e, and thus further supporting the jog-drag mechanism.\u003c/p\u003e\u003cp\u003eLater stages of extended yielding, specifically as plastic strain accumulates without increases in stress, can be explained by nucleation of dislocations and evolution of the linear complexion structures. Because solute-decorated dislocations are stronger sinks for irradiation-induced defects, such defects will accumulate more readily near decorated dislocations, and thus continuously increase the stress built up around them. These solute decorated dislocations containing a large concentration of defects, remain immobile as they require significantly higher activation stress than regular dislocation motion\u003csup\u003e\u003cspan citationid=\"CR75\" class=\"CitationRef\"\u003e75\u003c/span\u003e\u003c/sup\u003e. Hence, the excess stress experienced by these decorated dislocations leads to rapid and significant dislocation multiplication and subsequent motion of newly formed dislocations\u003csup\u003e\u003cspan citationid=\"CR37\" class=\"CitationRef\"\u003e37\u003c/span\u003e\u003c/sup\u003e, resulting in a dislocation avalanche and strain burst as shown by discrete dislocation dynamics simulation\u003csup\u003e\u003cspan citationid=\"CR76\" class=\"CitationRef\"\u003e76\u003c/span\u003e\u003c/sup\u003e, thus enabling dislocation plasticity to proceed without more need of bulk applied stress. Another possibility is that the structure of linear complexion changes during deformation, specifically that β-FeSi\u003csub\u003e2\u003c/sub\u003e may transform to α-FeSi\u003csub\u003e2\u003c/sub\u003e. The β-FeSi\u003csub\u003e2\u003c/sub\u003e\u0026rarr;α-FeSi\u003csub\u003e2\u003c/sub\u003e transformation can occur through either diffusion or deformation, the latter pathway being more energetically favorable\u003csup\u003e\u003cspan citationid=\"CR67\" class=\"CitationRef\"\u003e67\u003c/span\u003e\u003c/sup\u003e. The transformation energy for β-FeSi\u003csub\u003e2\u003c/sub\u003e\u0026rarr;α-FeSi\u003csub\u003e2\u003c/sub\u003e transformation has a local minimum at 0.5 strain, when the β-FeSi\u003csub\u003e2\u003c/sub\u003e structure is under monoclinic (100)[001] strain as shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig7\" class=\"InternalRef\"\u003e7\u003c/span\u003ec. However, since α-FeSi\u003csub\u003e2\u003c/sub\u003e is not stable at room temperature\u003csup\u003e\u003cspan citationid=\"CR57\" class=\"CitationRef\"\u003e57\u003c/span\u003e\u003c/sup\u003e, the α-FeSi\u003csub\u003e2\u003c/sub\u003e phase within dislocation segments is metastable and will ultimately compromise the ability to form or maintain linear complexion configurations. Thus, the combination of dislocation multiplication from decorated dislocations and β-FeSi\u003csub\u003e2\u003c/sub\u003e\u0026rarr;α-FeSi\u003csub\u003e2\u003c/sub\u003e structural changes inside linear complexions results in an end to linear complexion-mediated extended yielding.\u003c/p\u003e\u003cp\u003e\u003c/p\u003e\u003cp\u003eTo conclude, this study has documented the irradiation-induced formation of linear complexions in Grade 91 ferritic steel, whereas linear complexions have previously only been documented in model alloy systems. The β-FeSi\u003csub\u003e2\u003c/sub\u003e complexion formation is made possible because the lower thermal conductivity of the powder-based (PM-HIP) Grade 91 steel enables retention of a high density of screw dislocations during alloy fabrication and irradiation. The observed β-FeSi\u003csub\u003e2\u003c/sub\u003e complexions provide unprecedented improvements in ductility after irradiation, through screw dislocation-mediated plasticity. The formation of these linear complexions enables the alloy to overcome the strength\u0026ndash;ductility tradeoff, one of the most longstanding challenges that limit the service lifetime of nuclear structural alloys. These findings open new strategies for improving the mechanical performance of ferritic steels and other structural alloys commonly used in nuclear energy applications. More broadly, this work presents new directions for complexion engineering of future metallic alloys that are autonomously resilient to irradiation and offer unprecedented combinations of strength and ductility under irradiation.\u003c/p\u003e"},{"header":"Methods","content":"\u003cp\u003e\u003cb\u003eMaterials and irradiation\u003c/b\u003e. The materials selected for the current study are Grade 91 ferritic steels produced either by conventional casting or by powder metallurgy with hot isostatic pressing (PM-HIP) processes. The cast ingot and PM-HIP compact were provided by the Electric Power Research Institute (EPRI), USA. The PM-HIP was fabricated by consolidating gas atomized Grade 91 alloy powders under isostatic pressure of 103.4 MPa at a temperature of 1121\u0026deg;C for 4 hours. This was followed by a normalization heat treatment at 1060\u0026thinsp;\u0026plusmn;\u0026thinsp;14\u0026deg;C for 2.5 hours, with subsequent forced air fan cooling. Afterward, it was tempered at 777\u0026thinsp;\u0026plusmn;\u0026thinsp;14\u0026deg;C for 4.5 hours followed by air cooling. The cast ingot also underwent the same normalization and tempering treatments. The chemical compositions of the alloys were measured by inductively coupled plasma atomic emission spectroscopy (ICP-AES) and provided in \u003cb\u003eTable \u003cspan refid=\"MOESM1\" class=\"InternalRef\"\u003eS1\u003c/span\u003e.\u003c/b\u003e Chemical analysis confirms the alloys conform to Grade 91 steel specifications in ASTM A387, and have compositions consistent with other Grade 91 steel process heats studied specifically for nuclear applications\u003csup\u003e\u003cspan citationid=\"CR46\" class=\"CitationRef\"\u003e46\u003c/span\u003e\u003c/sup\u003e.\u003c/p\u003e\u003cp\u003eBoth the PM-HIP and cast materials were further sectioned into circular discs and round tensile bars. Disc specimens had a diameter of 3 mm and a thickness of 0.15 mm, and were prepared using wire electrical discharge machining (EDM). Discs were then mechanically polished using SiC paper up to 1200 grit, followed by electropolishing in a solution of 10% perchloric acid and 90% methanol (by volume) at \u0026minus;\u0026thinsp;40\u0026deg;C and 35 V for 20 seconds. Round threaded tensile bars were prepared to ASTM E8 specifications having dimensions of 76.2 mm in length and 6.35 mm in gauge diameter. The tensile bars were machined at Idaho National Laboratory (INL) using a computer numerical control (CNC) machining to a surface roughness of 3.2 \u0026micro;m. Comprehensive details of specimen machining and preparation are provided in ref.\u003csup\u003e\u003cspan citationid=\"CR77\" class=\"CitationRef\"\u003e77\u003c/span\u003e\u003c/sup\u003e.\u003c/p\u003e\u003cp\u003eFor neutron irradiations, tensile bar and disc specimens were loaded into drop-in capsules, which act as a barrier between the primary reactor coolant and the specimens, and were located in the A-6, A-7, and A-8 positions in the Advanced Test Reactor (ATR) at INL. Comprehensive details of the capsule design, irradiation thermal and fluence calculations and measurements, and irradiation histories, are provided in ref.\u003csup\u003e\u003cspan citationid=\"CR77\" class=\"CitationRef\"\u003e77\u003c/span\u003e\u003c/sup\u003e. A total of four irradiated disc specimens and six irradiated tensile bars were selected for microstructural and mechanical characterization, with additional unirradiated specimens reserved as controls. The target irradiation temperature was 400\u0026ordm;C and target irradiation doses were 1 displacement per atom (dpa) and 3 dpa. Actual irradiation doses and temperatures for each specimen are listed in \u003cb\u003eTable S2\u003c/b\u003e. For the 1 dpa target specimens, the average dose was 1.00\u0026thinsp;\u0026plusmn;\u0026thinsp;0.02 dpa at temperatures ranging 358\u0026ndash;389\u0026deg;C, whilst for the 3 dpa target specimens, the average dose of 3.68\u0026thinsp;\u0026plusmn;\u0026thinsp;0.23 dpa was achieved at temperatures 351\u0026ndash;397\u0026deg;C.\u003c/p\u003e\u003cp\u003e\u003cb\u003eMechanical testing and microstructural characterization.\u003c/b\u003e Quasi-static uniaxial tensile tests were conducted on irradiated tensile bar specimens, and on unirradiated reference specimens, following ASTM E8. The tests were performed at room temperature within an inert argon (Ar) atmosphere using a remote-operated Instron 5869 screw-driven load frame located inside the hot cells at the Hot Fuel Examination Facility (HFEF) Main Cell Window 13 M at the Materials and Fuels Complex (MFC), INL. Tensile tests were performed at an initial crosshead speed of 0.279 mm/min, corresponding to a strain rate of 8.78 \u0026times; 10\u003csup\u003e\u0026minus;\u0026thinsp;3\u003c/sup\u003e s\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e, until 10% engineering strain. Thereafter, crosshead speed was increased to 1.0 mm/min, i.e., strain rate of 3.15 \u0026times; 10\u003csup\u003e\u0026minus;\u0026thinsp;2\u003c/sup\u003e s\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e, until fracture.\u003c/p\u003e\u003cp\u003eMicrostructural characterization for the unirradiated and irradiated specimens was carried out on the disc specimens. Post-irradiation, the same electropolishing as was done prior to irradiation, was performed to remove minor surface damage resulting from handling, inspection, and decontamination. Focused ion beam (FIB) lift-out was conducted using a Thermo-Fisher Scientific (formerly FEI) Quanta 3D dual-beam scanning electron microscope (SEM), to extract lamellae for transmission electron microscopy (TEM) analyses, following established protocols\u003csup\u003e\u003cspan citationid=\"CR78\" class=\"CitationRef\"\u003e78\u003c/span\u003e\u003c/sup\u003e. To minimize FIB-induced damage, a protective Pt coating was deposited on the specimen surface prior to milling, mitigating Ga⁺ ion implantation and sputtering effects. A final polishing step using a 5 kV Ga⁺ ion beam, followed by 2 kV low-energy cleaning, was employed to mitigate surface damage and Ga⁺ implantation.\u003c/p\u003e\u003cp\u003eTEM observations were performed using a Thermo-Fisher Scientific Tecnai TF30-FEG STwin TEM, operated at 300 kV, and equipped with a high-angle annular dark field (HAADF) detector for scanning transmission electron microscopy (STEM), located at the Microscopy and Characterization Suite (MaCS) at the Center for Advanced Energy Studies (CAES). Dislocations and dislocation loops were imaged using the down-zone bright field (BF) STEM technique. This technique facilitates dislocation loop imaging by relaxing diffraction conditions\u0026mdash;specifically the g\u0026bull;b invisibility criterion\u0026mdash;through the use of a convergent electron beam. This allows simultaneous visualization of loops with different Burgers vectors and orientations, significantly simplifying their quantification and enhancing imaging efficiency. Compared to traditional two-beam or weak-beam dark-field imaging, BF-STEM offers higher contrast, reduced acquisition time, and improved statistical accuracy\u003csup\u003e\u003cspan citationid=\"CR79\" class=\"CitationRef\"\u003e79\u003c/span\u003e\u003c/sup\u003e. Void characterization was conducted using through-focus BF TEM, in which voids appear as contrast-reversing features between under- and over-focused images. The average lamella thickness used for volumetric density calculations was determined using electron energy loss spectroscopy (EELS) by analyzing the zero-loss peak in energy-filtered TEM mode. Image datasets were processed using TEM Instrument Analysis (TIA) software and measurements were done with ImageJ software.\u003c/p\u003e\u003cp\u003eThe irradiated disc specimens were further used for atom probe tomography (APT) studies. APT needles were prepared using FIB lift-out and annular milling, on the same Thermo-Fisher Scientific Quanta 3D FIB/SEM at MaCS, CAES. Local electrode atom probe (LEAP) analyses were carried out using a Cameca LEAP 4000X HR at MaCS, CAES, operated in laser-pulsed mode at a base temperature of 45\u0026ndash;50 K, 200 kHz pulse frequency, and 60 pJ laser energy. To ensure statistical confidence and reduce the influence of local heterogeneities in specimens, at least two needles per irradiation conditions were analyzed that exceeded 2\u0026nbsp;million ion detections per needle. APT data were reconstructed and analyzed using Cameca IVAS version 3.8.6 and AP Suite software.\u003c/p\u003e"},{"header":"Declarations","content":"\u003ch2\u003eAcknowledgements\u003c/h2\u003e\u003cp\u003eThe authors thank Dr. Donna Guillen, Jana Howard, and Jeremy Burgener of Idaho National Laboratory for their assistance with irradiation experiments and irradiated specimen handling; Dr. Zhongxia Shang of Purdue University for assistance with microscopy of unirradiated materials; Sri Sowmya Panuganti and Saquib Bin Habib of Purdue University for their assistance with data management; and Dr. Benjamin Sutton and David Gandy of the Electric Power Research Institute for material supply. Irradiation experiments and post-irradiation examination were supported by the U.S. Department of Energy, Office of Nuclear Energy, through the Nuclear Science User Facilities (NSUF) contracts 15-8242 and 23-4703. This project was also supported by the Electric Power Research Institute contract 10015819.\u003c/p\u003e"},{"header":"References","content":"\u003col\u003e\n\u003cli\u003eLi, Z., Pradeep, K. G., Deng, Y., Raabe, D. \u0026amp; Tasan, C. C. Metastable high-entropy dual-phase alloys overcome the strength\u0026ndash;ductility trade-off. \u003cem\u003eNature\u003c/em\u003e \u003cstrong\u003e534\u003c/strong\u003e, 227\u0026ndash;230 (2016).\u003c/li\u003e\n\u003cli\u003eWei, Y. \u003cem\u003eet al.\u003c/em\u003e Evading the strength\u0026ndash;ductility trade-off dilemma in steel through gradient hierarchical nanotwins. \u003cem\u003eNat. Commun.\u003c/em\u003e \u003cstrong\u003e5\u003c/strong\u003e, 3580 (2014).\u003c/li\u003e\n\u003cli\u003eRitchie, R. O. The conflicts between strength and toughness. \u003cem\u003eNat. Mater.\u003c/em\u003e \u003cstrong\u003e10\u003c/strong\u003e, 817\u0026ndash;822 (2011).\u003c/li\u003e\n\u003cli\u003eZhang, Z. 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Mater.\u003c/em\u003e \u003cstrong\u003e495\u003c/strong\u003e, 20\u0026ndash;26 (2017).\u003c/li\u003e\n\u003c/ol\u003e"}],"fulltextSource":"","fullText":"","funders":[],"hasAdminPriorityOnWorkflow":false,"hasManuscriptDocX":true,"hasOptedInToPreprint":true,"hasPassedJournalQc":"","hasAnyPriority":true,"hideJournal":false,"highlight":"","institution":"","isAcceptedByJournal":false,"isAuthorSuppliedPdf":false,"isDeskRejected":"","isHiddenFromSearch":false,"isInQc":false,"isInWorkflow":false,"isPdf":false,"isPdfUpToDate":false,"isWithdrawnOrRetracted":false,"journal":{"display":true,"email":"[email protected]","identity":"nature-portfolio","isNatureJournal":true,"hasQc":false,"allowDirectSubmit":false,"externalIdentity":"","sideBox":"","snPcode":"","submissionUrl":"","title":"Nature Portfolio","twitterHandle":"","acdcEnabled":false,"dfaEnabled":false,"editorialSystem":"ejp","reportingPortfolio":"","inReviewEnabled":true,"inReviewRevisionsEnabled":false},"keywords":"","lastPublishedDoi":"10.21203/rs.3.rs-7419483/v1","lastPublishedDoiUrl":"https://doi.org/10.21203/rs.3.rs-7419483/v1","license":{"name":"CC BY 4.0","url":"https://creativecommons.org/licenses/by/4.0/"},"manuscriptAbstract":"\u003cp\u003eHerein, the \u003cem\u003eoperando\u003c/em\u003e formation of Si-enriched linear complexions during neutron enables Grade 91 ferritic steel to overcome the strength\u0026ndash;ductility tradeoff, one of the most critical life-limiting challenges facing nuclear structural alloys. Linear complexions are a distinct yet confined chemical and structural state at a dislocation, which are rarely reported in engineering alloys. Ferritic steels are amongst the most ubiquitous engineering alloys for current and future nuclear components, but they are susceptible to irradiation hardening and embrittlement. Here, exceptional ductility retention exceeding 90% of pre-irradiation levels is obtained in Grade 91 synthesized using powder metallurgy with hot isostatic pressing (PM-HIP). Powder processing artifacts promote a high density of screw dislocation arrays, on which β-FeSi\u003csub\u003e2\u003c/sub\u003e linear complexions form due to Si segregation during irradiation. Screw dislocation dipoles undergo pinning and unpinning on linear complexions, resulting in extended yielding and excellent ductility retention post-irradiation. These findings represent a significant advancement toward design of alloys and manufacturing processes that can autonomously self-regulate their microstructural resilience \u003cem\u003ein-operando\u003c/em\u003e during irradiation, enabling exceptional ductility rather than embrittlement.\u003c/p\u003e","manuscriptTitle":"Linear complexions enable unprecedented ductility retention in neutron irradiated ferritic steel","msid":"","msnumber":"","nonDraftVersions":[{"code":1,"date":"2025-09-04 07:04:44","doi":"10.21203/rs.3.rs-7419483/v1","editorialEvents":[],"status":"published","journal":{"display":true,"email":"[email protected]","identity":"nature-communications","isNatureJournal":true,"hasQc":false,"allowDirectSubmit":false,"externalIdentity":"NCOMMS","sideBox":"Learn more about [Nature Communications](http://www.nature.com/ncomms/)","snPcode":"","submissionUrl":"https://mts-ncomms.nature.com/","title":"Nature Communications","twitterHandle":"","acdcEnabled":true,"dfaEnabled":true,"editorialSystem":"ejp","reportingPortfolio":"Nature Communications","inReviewEnabled":true,"inReviewRevisionsEnabled":false}}],"origin":"","ownerIdentity":"2199c126-00b8-4e02-b97d-be8128f881a8","owner":[],"postedDate":"September 4th, 2025","published":true,"recentEditorialEvents":[],"rejectedJournal":[],"revision":"","amendment":"","status":"under-review","subjectAreas":[{"id":54174150,"name":"Physical sciences/Materials science/Structural materials/Metals and alloys"},{"id":54174151,"name":"Physical sciences/Materials science/Structural materials/Mechanical properties"}],"tags":[],"updatedAt":"2026-05-10T01:50:14+00:00","versionOfRecord":[],"versionCreatedAt":"2025-09-04 07:04:44","video":"","vorDoi":"","vorDoiUrl":"","workflowStages":[]},"version":"v1","identity":"rs-7419483","journalConfig":"researchsquare"},"__N_SSP":true},"page":"/article/[identity]/[[...version]]","query":{"redirect":"/article/rs-7419483","identity":"rs-7419483","version":["v1"]},"buildId":"8U1c8b4HqxoKbykW_rLl7","isFallback":false,"isExperimentalCompile":false,"dynamicIds":[84888],"gssp":true,"scriptLoader":[]}