Effects of proton irradiation on nanocluster precipitation in ferritic steel containing fcc alloying additions

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Acta Materialia 60 (2012) 3034–3046 www.elsevier.com/locate/actamat

Effects of proton irradiation on nanocluster precipitation in ferritic steel containing fcc alloying additions Z.W. Zhang a,b,c, C.T. Liu a,d,⇑, X-.L. Wang b, M.K. Miller e, D. Ma b, G. Chen c, J.R. Williams f, B.A. Chin a a Materials Research & Education Center, Auburn University, Auburn, AL 36849, USA Chemical & Engineering Materials Division, Oak Ridge National Laboratory, Oak Ridge, TN 37831, USA c Engineering Research Center of Materials Behavior and Design, Ministry of Education, Nanjing University of Science and Technology, Nanjing 210094, People’s Republic of China d Center for Advanced Structural Materials, College of Science and Engineering, City University of Hong Kong, Kowloon, Hong Kong e Materials Science and Technology Division, Oak Ridge National Laboratory, Oak Ridge, TN 37831, USA f Physics Department, Auburn University, AL 36849, USA b

Received 25 December 2011; received in revised form 1 February 2012; accepted 3 February 2012 Available online 7 March 2012

Abstract Newly developed precipitate-strengthened ferritic steels with and without pre-existing nanoscale precipitates were irradiated with 4 MeV protons to a dose of 5 mdpa at 50 °C and subsequently examined by nanoindentation and atom probe tomography. Irradiation-enhanced precipitation and coarsening of pre-existing nanoscale precipitates were observed. Cu partitions to the precipitate core along with a segregation of Ni, Al and Mn to the precipitate/matrix interface after both thermal aging and proton irradiation. Proton irradiation induces the precipitation reaction and coarsening of pre-existing nanoscale precipitates, and these results are similar to a thermal aging process. The precipitation and coarsening of nanoscale precipitates are responsible for the changes in hardness. The observation of the radiation-induced softening is essentially due to the coarsening of the pre-existing Cu-rich nanoscale precipitates. The implication of the precipitation on the embrittlement of reactor-pressure-vessel steels after irradiation is discussed. Ó 2012 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Proton irradiation; Cu-rich nanoscale precipitate; Ferritic steel; Atom probe tomography; Solute enrichment

1. Introduction The long-term stability of structural materials in highly irradiated environments is a critical issue for future generations of advanced reactors and energy-related technologies [1–3]. Upon neutron irradiation in nuclear plants, Cu-containing steels, specifically the welds in reactorpressure-vessel (RPV) steels, exhibit radiation-enhanced embrittlement because the radiation-enhanced generation and diffusion of point defects and solute atoms accelerates ⇑ Corresponding author at: Center for Advanced Structural Materials, College of Science and Engineering, City University of Hong Kong, Kowloon, Hong Kong. E-mail address: [email protected] (C.T. Liu).

the formation of high densities of nanoscale Cu-rich precipitates, which act as obstacles to dislocations [2,4]. Owing to the potential importance of this phenomenon, many series of experiments have been carried out to study the structure evolution of precipitates in binary Fe–Cu and related alloys before and after electron, ion (including protons) or neutron irradiation [5–19]. Previous studies on RPV steels from surveillance capsules and other irradiated areas suggest that the radiation-induced Cu-rich precipitates remain stable under further irradiation [20,21]. As such, it is the stabilized Cu-rich precipitates that are responsible for the enhanced hardness that leads to embrittlement in neutron-irradiated RPV steels, as indicated by the shift in the ductile-to-brittle transformation temperature (DBTT) [22,23].

1359-6454/$36.00 Ó 2012 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2012.02.008

Z.W. Zhang et al. / Acta Materialia 60 (2012) 3034–3046

On the other hand, several other types of steels strengthened by similar or slightly larger Cu-rich precipitates, such as precipitation-hardened (PH) stainless steels [24–26] and the high-strength low-alloy (HSLA) steels [27–35], exhibit high strength and excellent toughness without embrittlement. The major difference between these two case is that the Cu-rich precipitates in the RPV steels have a body-centered cubic (bcc) structure, whereas as the size of the precipitates increases, as in the PH and HSLA steels, they change into the equilibrium e-Cu face-centered cubic (fcc) crystal structure [10,14]. Consequently, the interaction mechanisms with dislocations and their hardening response may change. Recent studies on nanostructured ferritic alloys (a subclass of oxide dispersion strengthened (ODS) ferritic steels) that contain high densities of stable nanoclusters can manifest excellent radiation tolerance [36–38]. Thus, in order to improve and design materials with better radiation resistance it is essential to understand the stability of preexisting Cu-rich nanoclusters under irradiation and their effects on radiation-induced defects. It is well known that displacement cascades produced during irradiation can promote the loss of these nanoscale particles by either ballistic dissolution or coarsening due to the radiation-enhanced diffusion of solute atoms [7,36]. To date, however, no detailed studies on the radiation damage resistance of nanoscale precipitate-strengthened ferritic steels containing Cu-rich nanoscale precipitates prior to irradiation have been reported. In this study, a newly developed Cu-rich nanoscale precipitate-strengthened ferritic steel [35,39,40] was used. The alloy design and thermal treatment routes produced a condition comprising a high-strength matrix with a high density of uniformly distributed nanoscale Cu-rich precipitates. To understand the effect of nanoscale features on the radiation resistance and, eventually the design of new alloys with superior tolerance to radiation, this steel was proton-irradiated to determine the radiation response. In contrast to neutron irradiation, ion irradiation provides a rapid and flexible way to achieve high doses without the hazards induced by material activation. A significant amount of work has been performed to relate damages induced by ion and neutron irradiation. It has been shown that displacement cascades induced by ion irradiation are similar to those induced by neutron irradiation [41]. Clarifying the irradiation response of Cu-rich precipitates requires detailed characterization of nanoscale precipitates, including size, morphology and composition before and after irradiation. Atom probe tomography (APT) is uniquely qualified for this purpose, due to its high spatial resolution and excellent elemental identification capacity. More importantly, APT provides a direct three-dimensional reconstruction of all precipitates on an atomic scale, and the composition of both the precipitate and matrix can be determined accurately. These capabilities are beneficial to identify the ballistic dissolution and the coarsening, both of which can induce the loss of nanoscale precipitates during irradiation.

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To elucidate both the irradiation effects on the stability of pre-existing Cu-rich precipitates and the direct correlations between the nanostructure and mechanical properties, the precipitates were characterized by APT before and after irradiation. Nanoindentation was used to determine the mechanical properties. Of great interest is quantitative determination of the evolution of Cu-rich precipitates and solute distribution/segregation in the matrix, precipitates and precipitate–matrix (P–M) interfaces under proton-irradiated conditions. The stability of pre-existing Cu-rich nanoscale precipitates and the solute segregation were studied and the implication of precipitation of Cu-rich nanoscale precipitates on the embrittlement of RPV steels after irradiation is discussed. The results are expected to be vitally useful for the design of irradiation-resistant alloys. 2. Experiment 2.1. Materials The chemical composition of the steel used for proton irradiation is listed in Table 1. Details of the material preparation have been described elsewhere [35]. Briefly, the Cucontaining ferritic steel was arc-melted and drop-cast into 10 mm diameter rods. The as-cast specimen was homogenized at 1000 °C for 1 h in air, followed by water quenching. The as-quenched specimen was heated at 950 °C for 30 min and then hot rolled through several passes to a thickness reduction of 80%. After a final pass the specimen was held at 950 °C for 15 min and then water quenched to suppress precipitation. The hot-rolled conditions have fully bcc crystal structure (i.e. ferritite) as determined by neutron diffraction using the VULCAN beamline [39,42,43] at the Spallation Neutron Source (SNS) [44] at Oak Ridge National Laboratory (ORNL). The hot-rolled specimens were then encapsulated in stainless steel tubes and cold-rolled to a thickness reduction of 60%. The finial thickness of the specimen was 130 lm. The cold-rolled specimens were ground and mechanically polished to 50 lm with a final surface roughness of 0.05 lm. The polished specimens were aged at 500 °C in vacuum (2  102 Pa) for 1 and 10 h, respectively. These specimens were then used for the proton irradiation. 2.2. Irradiation conditions Protons were used as sources to irradiate the target samples because of their relative large penetration depth. The proton irradiation experiments were conducted in the tandem accelerator laboratory at the Leach Science Center at Auburn University using 4 MeV protons and under a Table 1 Nominal chemical composition of the sample investigated in this study. Elements

Mn

Cu

Ni

Al

B

Fe

wt.% at.%

1.5 1.5

2.5 2.2

4.0 3.8

1.0 2.1

0.03 0.15

Bal. Bal

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vacuum of 104 Pa. Irradiation was performed at room temperature (25 °C) to separate and identify the effects of irradiation and thermal aging at 500 °C, respectively. However, proton irradiation induced a small average temperature raise of 25 °C to 50 ± 5 °C. The irradiation parameters are reported in Table 2. The proton beam was rastered over a uniformly irradiated area approximately 50 mm in diameter. The proton irradiation used a beam current of 4 lA which can provide a dose rate of 1.25  1012 protons cm2 s1. The proton distribution and profile of radiation-induced displacement as generated using the program Stopping and Range of Ions in Matter (SRIM) [45] for 4 MeV protons are shown in Fig. 1. The dose was based on the exposure fluence and the SRIM displacement calculation with 4 MeV protons and the detailed full damage cascade calculation method [45]. As can be seen from Fig. 1a, for a 50 lm thick specimen, most of the protons fully penetrated the specimen. Thus, a nearly uniform dose through the entire thickness of the specimen can be obtained which also minimized the influence of proton implantation, as shown in Fig. 1b. The damage dose profile along the depth of the specimen is shown in Fig. 2. The proton ion dose in the samples examined was 5 mdpa, which is comparable to a typical 40 year end-of-life exposure in a Westinghouse style nuclear reactor (30 mdpa) [46]. 2.3. Testing and characterization Due to the sample size limit, it is not practical to perform Charpy tests. Instead, the mechanical properties (hardness and Young’s modulus) of the specimens before

Fig. 2. Damage dose profile induced by proton irradiation along the specimen depth, showing the nearly uniform damage through the specimen.

and after proton irradiation were evaluated by nanoindentation in depth-control mode under ambient atmosphere with a MTS Nanoindenter XP system. The nanoindenter tip was a Berkovich-type diamond pyramid. The allowable thermal drift rate was limited to 0.05 nm s1. The continuous stiffness measurement (CSM) method was applied and the indentation depth was controlled within 2000 nm in order to evaluate the effect of irradiation on the mechanical properties at depth [47]. The Young’s modulus (E) and hardness (H) values were derived from the load–displacement curves at the beginning of the unloading segment using the Oliver–Pharr method [48]. For each sample, 16 measurements were taken to obtain the average values of E and H. To identify the radiation response of Cu-rich precipitates to proton irradiation, control samples and aged samples were selected to conduct the APT observation before

Table 2 Proton irradiation conditions for Cu-rich precipitate-strengthened ferritic steel. Ion Proton

Energy (MeV) 4

Dose rate (ion cm2 s1) 12

1.25  10

Fluence (ion cm2) 16

6.65  10

Total exposure (mdpa)

Temperature (°C)

Vacuum (Pa)

5

50 ± 5

1  104

Fig. 1. Results obtained from SRIM 2008 calculations for proton irradiation with 4 MeV protons showing: (a) the proton range distribution during irradiation and (b) the target displacement profile after irradiation.

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and after proton irradiation. The specimens without aging and radiation were used as control sample (CS) and the aged samples were aged for 10 h at 500 °C (AG10). The irradiated CS and AG10 specimens were labeled as CSIR and AG10IR conditions, respectively. APT specimens were prepared by a lift-out method in FEI Nova 200 focused ion beam (FIB) instrument. Typically, a 15 lm  3 lm  3 lm blank was cut from the samples and then attached to a tungsten micropost by Pt deposit. A series of annular milling operations was used to obtain sharp tips suitable for APT. APT was performed in voltage-pulsed mode with a Cameca Instruments (formerly Imago Scientific Instruments) local electrode atomprobe (LEAP 4000X HR). A specimen temperature of 50 K, a pulse repetition rate of 200 kHz and a pulse fraction of 0.2 were used for the analyses. This high pulse repetition rate significantly reduces the possibility of preferential evaporation of the low evaporation field solutes, such as Cu [49]. Imago Visualization and Analysis Software (IVAS) version 3.6 was used for 3-D reconstruction, composition analyses and to create isoconcentration surfaces. In addition to the atom maps and isoconcentration surfaces, the Cu-rich precipitates were identified with the maximum separation method [50], with a maximum separation distance of 0.5 nm, a minimum of 20 Cu atoms in the precipitates and a grid resolution of 0.1 nm. This maximum separation, or friends-of-friends, method enables the solutes in the precipitates to be distinguished from the solutes in the matrix so that their size, number density and concentration can be estimated. The size of the solute-enriched features was estimated in terms of the radius of gyration lg, which is determined from positions of the solute atoms in each precipitate. The radius of gyration was then converted to Guinier radius RG which presents the actual size of the feapffiffiffiffiffiffiffiffiffiffiffi tures according to equation: RG ¼ 5=3lg . Precipitate number densities, Nv, are calculated from [51]: Nv ¼

N pf ; nX

ð1Þ

where Np is the number of precipitates in the analyzed volume, n the total number of atoms detected in the same volume, f the detection efficiency of the single atom position-sensitive detector and X the average atomic volume. The solute distributions, associated with the precipitates, were generated with the proximity histogram method [52]. This method uses an isoconcentration surface to define a precipitate–matrix interface, and then measures the concentrations in discrete shells at fixed distances from that isoconcentration surface. The matrix composition of the each condition was also determined.

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shown in Figs. 3 and 4, respectively. The Young’s modulus remains relatively constant with various heat treatments. Proton irradiation does not induce a significant change in Young’s modulus. However, both heat treatment and proton irradiation induce a significant change in hardness, as shown in Fig. 4. For the specimen without proton irradiation, the hardness increases significantly after aging at 500 °C for 1 h due to the precipitation of Cu-rich precipitates. After aging for 10 h, the hardness increases further to 6 GPa, more than 50% increase in hardness, as compared with that of the CS condition. The hardness of the CS sample increased significantly from 3.9 to 5.6 GPa after proton irradiation, while proton irradiation only produced a slight increase in hardness of the sample aged for 1 h at 500 °C, as shown in Fig. 4. The hardness of the samples aged for 1 h at 500 °C and unaged exhibited a similar value of 5.6 GPa after the proton irradiation. These results indicate that the pre-existing precipitates with an optimum size induced by aging can alleviate the effect of proton irradiation on the hardness. For the sample aged for 10 h, the hardness decreased after irradiation, indicating a radiation-induced softening. To understand these points, APT needs to be used to study the precipitation evolution.

Fig. 3. Young’s modulus of CS, AG1 and AG10 conditions with different heat treatments and before and after proton irradiation. CS is the control sample without aging and radiation; AG1 and AG10 are the samples aged for 1 and 10 h at 500 °C, respectively.

3. Results 3.1. Mechanical properties The Young’s modulus and hardness, as determined by nanoindentation before and after proton irradiation, are

Fig. 4. Hardness of CS, AG1 and AG10 conditions before and after proton irradiation. CS is the control sample without aging and radiation; AG1 and AG10 are the samples aged for 1 and 10 h at 500 °C, respectively.

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3.2. Evolution of the precipitate properties The 3-D atom maps of the CS and CSIR conditions obtained by APT are shown in Fig. 5. As expected, no precipitates were present in the unirradiated control sample. All the solute elements (Cu, Ni, Al and Mn) were distributed uniformly in a supersaturated solid solution (Fig. 5a). However, after irradiation, a high number density of small precipitates was present, indicating that proton irradiation results in the precipitation of Cu-rich precipitates enriched in Ni, Al and Mn. The Ni and Mn enrichments are commonly observed in RPV steels [6,7,19,53]. The solute distributions in the AG10 condition before and after irradiation are shown in the atom maps in Fig. 6. Large numbers of precipitates, which were enriched in Cu, Ni, Al and Mn, were observed in both unirradiated and proton-irradiated conditions. This is consistent with the results obtained by Kolli et al. [54] in another Fe–Cu multicomponent steel after aging. Isoconcentration surfaces of Cu-rich precipitates in the CSIR, AG10 and AG10IR conditions are shown in Fig. 7. The isoconcentration surface is constructed at 16 at.% Cu. The well-defined morphology of precipitates is evident for all CSIR, AG10 and AG10IR conditions. The Cu-rich precipitate size distributions after different heat treatments and irradiation processes are shown in Fig. 8. The average Guinier radius, hRG i, and number density, Nv, are also included. The precipitate size distributions are well described by a log-normal function [55], as shown in Fig. 8. The fitting curves are overlaid onto the experimental data obtained by APT. The average Guinier radii of the

CSIR, AG10 and AG10IR conditions are 1.6, 1.7 and 2.1 nm, respectively, and the number densities are 1.8  1024, 1.1  1024 and 6.4  1023 m3, respectively. These values show that after irradiation the precipitate sizes of CSIR and AG10 conditions are similar. However, for the AG10 condition, the precipitate size increases significantly after irradiation. Correspondingly, the number density is reduced significantly (nearly 50%), indicating that precipitate coarsening occurred during irradiation. The polydispersity of the size distributions determined from log-normal fits of CSIR, AG10 and AG10IR conditions are 0.29, 0.32 and 0.34, respectively, indicating a slight broadening of size distribution which is similar to the aging process. In this APT study, no grain boundaries or dislocations were encountered due to the large grain size of the material. 3.3. Evolution of compositions The distribution of the solute elements Cu, Ni, Mn and Al are depicted clearly in the atom maps shown in Figs. 5 and 6. The enriched regions of Cu, Ni, Mn and Al atoms are indicative of ultrafine precipitates. Both irradiated and thermally aged conditions exhibit a similar portioning trend. Proximity histograms showing the solution partitioning to the precipitates and matrix are presented in Fig. 9 for Fe, Cu, Ni, Al and Mn for the CSIR, AG10 and AG10IR conditions. It is evident that the depletion of Fe and the enrichment of Cu change monotonically towards the center of the precipitates, whereas the concentration of Ni, Al and Mn exhibits a diffuse enrichment near the P–M interfaces.

Fig. 5. Three-dimensional APT atom maps of the solute distribution of Cu, Ni, Al and Mn elements in the CS samples: (a) before irradiation (CS condition) and (b) after irradiation (CSIR condition).

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Fig. 6. APT atom maps showing the distribution of Cu, Ni, Al and Mn solute elements in the samples aged for 10 h at 500 °C: (a) before (AG10 condition) and (b) after proton irradiation (AG10IR condition).

Fig. 7. Copper-rich precipitates as delineated by 16 at.% Cu isoconcentration surfaces in (a) the irradiated CS sample (CSIR), (b) the sample aged at 500 °C for 10 h before radiation (AG10) and (c) the sample aged at 500 °C for 10 h and irradiation (AG10IR).

Fig. 8. Cu-rich precipitate size distribution in CSIR, AG10 and AG10IR conditions, respectively. The superimposed curves are log-normal fits to the APT data. AG10, sample aged for 10 h at 500 °C; CSIR, irradiated CS condition; AG10IR, irradiated AG10 condition.

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Fig. 9. Proximity histograms (at.%) for Fe, Ni, Cu, Mn and Al determined for the (a) CSIR, (b) AG10 and (c) AG10IR conditions. The vertical dashed line in each graph corresponds to the isoconcentration level used to mark the P–M interface between matrix and Cu-rich precipitates. AG10, sample aged for 10 h at 500 °C; CSIR, irradiated CS condition; AG10IR, irradiated AG10 condition.

3.3.1. Composition of the precipitates The composition of precipitates in CSIR, AG10 and AG10IR conditions are summarized in Table 3. These compositions were estimated from central region of the precipitate with a radius of 0.5 nm based on information from the proximity histogram. In the CSIR condition, the precipitate cores are Cu-rich (60.6 ± 2.7 at.%) but also contain significant concentrations of Fe, Ni, Al and Mn. In the AG10 condition, the precipitate cores are also enriched in Cu (66.0 ± 3.1 at.%) with significant concentrations of Fe, Ni, Al and Mn. After irradiation of AG10 samples (i.e. AG10IR), the Cu concentration within the precipitate cores increases significantly to 83.4 ± 5.2 at.% with concomitant decreases in the concentration of Fe, Ni and Al as compared to the unirradiated AG10 condition. The concentrations decrease to 4.9 ± 1.4 at.% Ni, 4.0 ± 1.6 at.% Al and 4.2 ± 1.9 at.% Fe along with a slight increase in Mn to 3.5 ± 0.6 at.% Mn. These results may suggest that the precipitates are approaching their equilibrium concentrations after irradiation. 3.3.2. Composition of the a-Fe matrix To identify the evolution of nanoscale Cu-rich precipitates by either ballistic dissolution or coarsening under proton irradiation, the composition of matrix was determined before and after irradiation. For the CS condition, the esti-

mation of the matrix composition was not complicated by the presence of a high number density of precipitates, so the entire datasets were used for the estimates. In this case, the matrix composition can be compared to the bulk composition if it is assumed that there are no coarse precipitates or inclusions present in the microstructure. For the conditions containing precipitates, the precipitates were identified with a 3 at.% Cu isoconcentration surface and were then removed from the APT data. This threshold concentration was chosen as it is slightly higher than the bulk Cu value and far lower than the Cu concentration within the precipitates. These precipitate-free APT data were used to determine the composition of the a-Fe matrix and the results are listed in Table 3. As compared to the CS condition, the concentration of Fe within the matrix in the CSIR, AG10 and AG10IR conditions increases due to the formation of Cu-rich precipitates and ejection of Fe. Accordingly, the matrix concentration of Cu decreases significantly from 1.85 at.% to 0.35, 0.32 and 0.22 at.% in the CSIR, AG10 and AG10IR conditions, respectively. After irradiation, the Cu concentration in the AG10IR matrix decreases further as compared to the level before irradiation. This indicates that proton irradiation does not redistribute the Cu into the matrix. This result is consistent with the enrichment of Cu in precipitates under proton irradiation, as shown in Table 3. The concentrations of Ni and Al

Z.W. Zhang et al. / Acta Materialia 60 (2012) 3034–3046 Table 3 Compositions in the precipitates and matrix and at the P–M interfaces as determined by APT. Element

CS (at.%)

Precipitates Cu Ni Al Mn Fe Matrix Cu Ni Al Mn Fe Interfaces Cu Ni Al Mn Fe

1.85 ± 0.02 3.95 ± 0.03 3.04 ± 0.03 1.11 ± 0.01 90.05 ± 0.04

CSIR (at.%)

AG10 (at.%)

AG10IR (at.%)

60.6 ± 2.7 8.3 ± 1.5 8.3 ± 0.9 3.3 ± 0.8 19.5 ± 1.5

66.0 ± 3.1 9.4 ± 2.5 10.4 ± 2.3 2.7 ± 0.6 11.5 ± 2.5

83.4 ± 5.2 4.9 ± 1.4 4.0 ± 1.6 3.5 ± 0.6 4.2 ± 1.9

0.35 ± 0.003 3.18 ± 0.010 1.39 ± 0.007 1.05 ± 0.005 94.02 ± 0.01

0.32 ± 0.006 2.57 ± 0.014 1.33 ± 0.009 1.03 ± 0.009 94.75 ± 0.02

0.22 ± 0.006 2.64 ± 0.010 1.27 ± 0.009 1.00 ± 0.009 94.88 ± 0.02

16.7 ± 4.7 12.4 ± 0.8 8.3 ± 0.6 3.3 ± 0.2 59.3 ± 6.1

16.3 ± 3.2 16.1 ± 1.3 12.2 ± 1.2 2.8 ± 0.4 52.7 ± 4.7

16.4 ± 3.9 17.9 ± 1.2 10.2 ± 1.2 4.3 ± 0.3 51.0 ± 6.4

The composition of precipitates was estimated from the precipitate core with a radius of 0.5 nm based on the proximity histogram. The composition of interfaces was estimated from regions within ±0.5 nm of the interfaces from the proximity histogram. CS, control sample without aging and radiation; AG10, sample aged for 10 h at 500 °C; CSIR, irradiated CS condition; AG10IR, irradiated AG10 condition.

are also depleted in the matrix after irradiation and aging, as compared to the CS condition, which is consistent with the partitioning of these elements to the precipitates and the P–M interfaces. 3.3.3. Solute enrichment at the P–M interfaces The proximity histograms across the P–M interfaces show the enrichment of Ni, Al and Mn for the CSIR, AG10, AG10IR conditions, as shown in Fig. 9. The composition of interfaces was estimated from the proximity histogram from regions within ±0.5 nm of the interfaces, and the results are list in Table 3. For the CSIR, AG10 and AG10IR conditions, the Cu concentrations are comparable. However, irradiation induces increases in Ni and Mn from 16.1 at.% Ni and 2.8 at.% Mn in the AG10 condition to 17.9 at.% Ni and 4.3 at.% Mn in the AG10IR condition. The Al concentration at the interfaces decreases slightly

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from 12.2 at.% Al in the unirradiated AG10 condition to 10.2 at.% Al in AG10IR condition. The enrichment of Ni at the interfaces is detected for the majority of the Cu-rich precipitates in an a-Fe matrix [54,56–58]. These results are consistent with the Ni interfacial excesses derived from proximity histograms [54,56]. 4. Discussion 4.1. Radiation-induced precipitates vs. aging-induced precipitates The microhardness of the CS condition increases significantly from 3.9 to 5.2 GPa after proton irradiation (see Fig. 4). This increase can be related to the formation of Curich precipitates with a radius of 1.6 nm and a number density of 1.8  1024 m3, as revealed in the atom maps in Fig. 5b. They are obstacles to the movement of dislocations under stress [59]. Proton irradiation may lead to a different mechanism of nucleation from the thermal aging process, because radiation damage arising from the displacement cascades can form vacancies and self-interstitial atoms (SIAs), dislocation loops and in some cases small voids [2,3,60]. Therefore, a comparison of the mechanical properties, precipitate parameters and composition variations during irradiation and aging processes was performed. The hardnesses of the thermally aged AG10 and irradiated CSIR conditions are comparable, as shown in Fig. 4. This similarity correlates with the comparable sizes and number densities of Cu-rich precipitates in these two conditions (Fig. 8). Local composition variations at the enrichment site may influence the nucleation mechanism of precipitates between the irradiation and aging processes. Composition variation of the Cu-rich precipitates and P–M interfaces, as shown in Fig. 10, indicates that all the solute elements Cu, Ni, Al and Mn have similar partition patterns in the irradiated CSIR condition as in the aged AG10 condition. The proximity histogram concentration profiles shown in Fig. 9 provide additional evidence that in both irradiated CSIR and aged AG10 conditions, the Ni, Al and Mn elements are enriched at the interfaces. All these results indicate that the mechanism of precipitation under the proton irradiation is similar to that of the thermal aging.

Fig. 10. Concentrations of Fe, Cu, Ni, Al and Mn elements (a) in precipitates and (b) at P–M interfaces in irradiated CSIR and unirradiated AG10 conditions, respectively. AG10, sample aged for 10 h at 500 °C; CSIR, irradiated CS condition.

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The main effect of the proton irradiation is to enhance the diffusion of Cu due to the production of freely migrating vacancies [61]. The effect of irradiation on the formation of Cu-rich precipitates can be taken into account by using the radiation-enhanced diffusion coefficient of Cu atoms [8]: e DirCu ¼ Dth Cu C v =C v ;

ð2Þ

DirCu

where is radiation-enhanced diffusion coefficient of Cu atoms, and Dth Cu is the thermal diffusion coefficient of Cu atoms. Cv and C ev are the vacancy concentrations under irradiation and thermal-equilibrium conditions, respectively. The vacancy concentration is related to both the recombination reactions and loss of point defects at sinks as [8,61]: sffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi!1 2G 4Gl Dv C v ¼ 2 1 þ 1 þ 4 R ; ð3Þ k Dv k where G is the effective generation rate of Frenkel pairs, k is the sink strength for the point defects, lR is the recombination constant and Dv is the vacancy diffusion coefficient. The dose rate of proton irradiation is 9.4  108 dpa s1 in this study. Other parameters were taken as: lR = 1021 m2, k2 = 1014 m2 and eR = 0 [8]. The vacancy  diffu m

sion coefficient was taken to be Dv ¼ 5  105 exp  kEBvT   eV m2 s1 [8,62], and that for DV C eV ¼ 5  104 exp  2:77 kB T

m2 s1 [8,62], where kBT has the standard meaning. The vacancy diffusion activation energy Emv can be varied from 0.53 to 1.22 eV for pure Fe [8,63] and ferritic steels [8,62], respectively. Accordingly, the radiation-enhanced diffusion coefficient, DirCu , determined by Eq. (2) can be varied from 9.7  1019 down to 1.2  1023 m2 s1. According to diffusion experiments by Anand and Agarwala [64], the Cu diffusion coefficient under thermal aging at 500 °C can be 21 m2 s1 . Assuming that this estimated as Dth Cu ¼ 1:7  10 Cu thermal diffusion coefficient is similar to the radiation-enhanced diffusion coefficient at 50 °C, a vacancy diffusion activation energy of Emv ¼ 0:94 eV, can be obtained which is reasonable when taking into account the alloy composition and the higher dislocation density resulting from cold-rolling before proton irradiation. This result supports the mechanism that the point defects induced by high-energy cascades during irradiation may enhance the precipitation kinetics [60,65]. It is reported that for an Fe–1.34 at.% Cu alloy under electron irradiation, as well as thermal aging, precipitate kinetic data can be explained only by assuming that small embryos are present at the beginning with a number density independent of the temperature [66,67]. In our recent study [39], by using combined small-angle neutron scattering (SANS) and APT experiments, embryos were found to exist with a size of 0.4 nm in an as-quenched sample due to sub-nanoscale density fluctuation or phase separation. Upon irradiation, as well as annealing, these local regions of Cu enrichment may be the sites at which the precipitates

can nucleate and grow. In addition, Monte Carlo simulations of Cu-rich precipitates show that small Cu-rich complexes can also contribute significantly to the transport of Cu atoms and may contribute to the formation of Cu-rich precipitates [68]. In this study, proton irradiation was found to result in similar precipitate sizes and number densities in both the CSIR and AG10 conditions. Moreover, there are similar enrichments of Cu, Ni and Al atoms, as shown in Fig. 10. All these results provide direct evidence that both proton irradiation and thermal aging have similar effects on the kinetics of precipitate growth. Proton irradiation enhances the diffusion of solute elements due to the production of freely migrating point defects similar to solute diffusion under thermal aging. 4.2. Effect of irradiation on pre-existing Cu-rich precipitates 4.2.1. Effect of irradiation on the composition evolution of pre-existing Cu-rich precipitates By comparing the precipitate compositions before and after irradiation of the AG10 condition (see Fig. 11), it is evident that the irradiation process results in a change in their compositions. The concentration of Cu increased significantly in the precipitate and the concentrations of Fe, Ni and Al decreased significantly. The Cu concentration at interfaces did not change. The Al concentration decreased and the Ni and Mn concentrations increased. Isheim et al. [69] and Kolli et al. [54,70] have reported similar results under thermal aging in related Fe–Cu-based steels. They suggested that Ni and Al atoms can form the NiAl-type B2-ordered phase with Mn substituting for Al, forming a Ni–Al–Mn shell around a Cu-rich core after prolonged aging. First-principles calculations also confirm that Mn substitutes on the Al sublattice sites rather than on the Ni sublattice sites. In this study, after aging for 10 h at 500 °C, the concentration of Cu in precipitates was 66.0 ± 3.1 at.% and proton irradiation was found to significantly increase the Cu concentration in the precipitates to 83.4 ± 5.2 at.% along with decreases of Ni, Al and Mn concentrations. This indicates that irradiation induces the evolution of pre-existing Cu-rich precipitates approaching a steady-state condition from a metastable state. By comparing the composition evolution after prolonged thermal aging and after irradiation, low-temperature proton irradiation may play a similar role to thermal aging in the evolution of pre-existing Cu-rich precipitates. 4.2.2. Effect of irradiation on the stability of pre-existing Curich precipitates According to theory of precipitate stability under irradiation presented by Russell [71], precipitates can grow through diffusion of solutes or solute clusters in the matrix or can shrink due to radiation-induced recoils and ballistic ejection of precipitates. The ballistic ejection of atoms from precipitates in the matrix has several possible outcomes, such as back-diffusion to rejoin the original particle, diffusion to join another particle, help nucleate a new particle,

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Fig. 11. Concentrations of Fe, Cu, Ni, Al and Mn elements (a) in precipitates and (b) at the P–M interfaces in AG10 and AG10IR conditions. AG10, sample aged for 10 h at 500 °C; AG10IR, the irradiated AG10 condition.

or remain in solution in the matrix. All these cases can influence the stability of Cu-rich precipitates. As can be seen from Figs. 7 and 8, after irradiation of the AG10 condition, the precipitate size increased from 1.7 to 2.1 nm. Correspondingly, the number density decreased from 11  1023 to 6.4  1023 m3 after irradiation, a 50% decrease. A slight decrease in the concentration of Cu in the matrix from 0.32 to 0.22 at.% was also observed after irradiation (Table 3). These results demonstrate that the precipitates did not dissolve into the matrix but coarsened during irradiation. This coarsening is responsible for the decrease in hardness after irradiation. 4.3. Effect of pre-existing Cu-rich precipitates on irradiation tolerance Many studies have observed radiation-enhanced embrittlement, especially in Cu-containing RPV steels and related alloys [53,72–75]. Radiation-enhanced embrittlement is generally attributed to the presence of high number densities of Cu-, Ni-, Si- and Mn-enriched precipitates, which hinder the movement of dislocations, although there is only small a amount of Cu in these kinds of steels, typically less than 0.15 wt.% Cu in solid solution after the post-weld heat treatment. On the other hand, the Cu-rich precipitatestrengthened steels with a high concentration of Cu, typically higher than 1 wt.% Cu, present excellent mechanical properties with high strength and ductility [27–29]. In the present study, radiation-induced softening in AG10, which contains pre-existing Cu-rich nanoscale precipitates, was observed for the first time. This clearly indicates that the pre-existing Cu-rich precipitates are beneficial to the softening of steels under irradiation. Despite the microstructural differences between RPV steels and the alloy used in this investigation, a comparison can be made. The Cu-rich/enriched precipitates have been characterized in many neutron-irradiated RPV and related steels by atom probe field ion microscopy (APFIM) and APT [6,7,76–79]. Neutron irradiation can induce a high number density of Cu-rich precipitates and a decrease in DBTT. For example, neutron irradiation of an A533B steel [49] to a fluence of 5  1023 nm2 resulted in a high number

density of 3  1023 m3 Cu-enriched precipitates with an average Guinier radius of 1.4 nm. The average composition of the precipitates was 87 at.% Fe along with significant enrichment of Ni, Mn and Si in the precipitates. After neutron irradiation, the DBTT increased by 96 °C [49], indicating radiation-enhanced embrittlement. After a post-irradiation annealing treatment for 168 h at 460 °C, the number density of the Cu-rich precipitates decreased to 2  1022 m3 and their size coarsened to 1.9 nm [49]. In addition, the DBTT was fully recovered. Similar phenomena were also observed in a in-service irradiated French RPV steel (Chooz A) [77]. This is similar to the proton irradiation results from the AG10 condition in which the sample containing pre-existing Cu-rich precipitates was softened after irradiation due to radiation-induced Cu-rich precipitate coarsening (Figs. 4 and 8). These results indicate that the existence of Cu-rich precipitates is not the dominant factor inducing embrittlement. Pareige et al. [80] observed by APT that a fluence of 9.7  1023 nm2 (E > 0.5 MeV) neutron irradiation induced Cu-enriched precipitates with a precipitate size and number density of 0.8 nm and 5  1023 m3, respectively in 15Kh2MFA type RPV steels [81]. The Cu concentration in the precipitates is 35 at.% along with significant amounts of Ni, Mn, Cr and Si. This low Cu concentration is consistent with the results obtained by thermal aging when the cluster size is smaller in Cu-containing ferritic steel [54]. Re-irradiation of the 15Kh2MFA steel after primary irradiation and stress relaxation thermal annealing resulted in a coarsening of Cu-rich precipitates to 2.5 nm. The Cu concentration increased to 81 at.% after re-irradiation, and re-irradiation did not promote a formation of new Cu-rich clusters [80]. These results are very similar to those in our proton-irradiated AG10 condition, AG10IR, where irradiation induced a coarsening of pre-existing Cu-rich precipitates (Fig. 8 and Table 3). However, re-irradiation leads to a re-embrittlement in the RPV steel, whereas in our proton-irradiated condition, AG10IR, the hardness decreased after proton irradiation. The re-embrittlement in the RPV steel was attributed to the segregation of impurities because re-irradiation promoted phosphorus-enriched zones and carbide segregation on the dislocations and grain

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boundaries [80]. All these studies of RPV steels are further complicated by the segregation of P to the dislocations which will have a strong influence on the dislocation–precipitate interaction and hence the embrittlement of these alloys. Carter et al. [82] and Nishiyama et al. [73] also observed that neutron irradiation promoted intergranular P segregation, which more significantly affects radiationenhanced hardening than that of Cu. The segregation of P is not present in the steel used in this study, as no impurity segregation was observed. However, in our very recent studies on the Cu-rich precipitate-strengthened ferritic steels [40], it was found that the hardened steels are very sensitive to impurities and introduced oxygen and hydrogen. The well-controlled Cu-rich precipitate-strengthened samples produced by vacuum heat treatment demonstrated excellent mechanical properties with high yield strength and satisfactory ductility, whereas hardened samples containing even small amounts of oxygen or hydrogen can be very brittle. Putting all these results together, the scenario is that the hardness can be increased significantly along with satisfactory toughness and ductility by the formation of Cu-rich precipitates, as demonstrated in PH and HSLA steels. However, these hardened steels may be very sensitive to impurity segregation and irradiation-induced point defects, such as SIAs, vacancies, voids, helium bubbles, etc., which emerge through displacement cascades and transmutation under neutron irradiation [83–85]. Cu precipitates in Fe can be sinks for both vacancy and interstitial point defects, and hence can act as recombination centers under irradiation conditions [22,86]. Therefore, the pre-existing Cu-rich precipitates are a promising approach to enhance the radiation resistance of steels that are in a different operating regime compared to RPV steels. 5. Conclusions A newly developed Cu-rich nanoscale precipitatestrengthened ferritic steel prepared under different conditions was irradiated by a proton source. By controlling the sample thickness and proton source energy, a uniform damage dose and minimized ion-implantation effect, both of which are similar to the neutron irradiation conditions in a reactor, were obtained. Radiation-induced precipitation and the stability of pre-existing Cu-rich nanoscale precipitates under proton irradiation and their effects on the mechanical properties were carefully examined by nanoindentation and APT. The following important findings were found: 1. There are no effects of proton irradiation and aging process on the Young’s modulus, whereas the hardness increases significantly after aging and proton irradiation of solution-treated CS due to the formation of Cu-rich precipitates.

2. Unusual irradiation-induced softening was observed after proton irradiation in the condition containing pre-existing Cu-rich nanoscale precipitates. 3. Irradiating at 50 °C to 5 mdpa leads to a high density of Cu-rich nanoscale precipitates in the CSIR condition. The Guinier radius and number density of Cu-rich precipitates is 1.6 nm and 1.8  1024 m3, respectively. The concentration of Cu is 61 at.% with significant levels of Fe, Ni, Al and Mn elements. The precipitate size, number density and composition of the precipitates are all comparable with that after aging at 500 °C for 10 h, i.e. the AG10 condition. 4. After irradiation of the AG10 condition, the Guinier radius increases from 1.7 to 2.1 nm and the number density decreases from 11  1023 to 6.4  1023 m3. In addition, the Cu matrix concentration decreased. All these results indicate that irradiation does not induce dissolution of precipitates but rather results in coarsening similar to the effect of thermal aging. This contributes to the radiation-induced softening. 5. Under thermal aging and irradiation, Cu partitions to the precipitate whereas Ni, Al and Mn enrich mainly at the P–M interface, thereby forming a Ni–Al–Mnenriched shell. 6. The Cu precipitation and the segregation of Mn at the P–M interface were enhanced by irradiation. The mechanisms of radiation-induced softening and radiation-enhanced embrittlement have been discussed, based on atomistic studies of how the nanoscale Cu-rich precipitates evolve under irradiation and their effects on the mechanical properties. This observed nanostructural evolution is expected to provide unique and important information about novel alloy design to enhance radiation resistance. Acknowledgements We thank the Tandem Accelerator Laboratory at the Leach Science Center at Auburn University for the irradiation instrument time and we acknowledge the help of our local contacts Tamara Isaacs-Smith and Max Cichon. This research was supported by internal funding from Auburn University and City University of Hong Kong, together with the NJUST Research Funding (No. 2010GIPY031), the NSFC Funding (Nos. 50871054 and 51171081) and RFDP Funding (No. 20113219120044). M.K.M., X.L.W. and Z.W.Z. were sponsored in part by the Division of Materials Science and Engineering, Office of Basic Energy Sciences, US Department of Energy. Atom probe tomography research (M.K.M., Z.W.Z.) at the Oak Ridge National Laboratory SHaRE User Facility was sponsored by the Scientific User Facilities Division, Office of Basic Energy Sciences, US Department of Energy.

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