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Deuterium ion irradiation induced precipitation in Fe–Cr alloy: Characterization and effects on irradiation behavior
Journal of Nuclear Materials 459 (2015) 81–89
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Journal of Nuclear Materials journal homepage: www.elsevier.com/locate/jnucmat
Deuterium ion irradiation induced precipitation in Fe–Cr alloy: Characterization and effects on irradiation behavior P.P. Liu a, R. Yu b, Y.M. Zhu a, M.Z. Zhao a, J.W. Bai a, F.R. Wan a, Q. Zhan a,⇑ a b
School of Materials Science and Engineering, University of Science and Technology Beijing, Beijing 100083, China Department of Materials Science and Engineering, Beijing National Center for Electron Microscopy, Tsinghua University, Beijing 100084, China
h i g h l i g h t s A new phase precipitated in Fe–Cr alloy after deuterium ion irradiation at 773 K. B2 structure was proposed for the Cr-rich new phase. Strain fields around the precipitate have been measured by GPA. The precipitate decrease growth rate of dislocation loop under electron irradiation.
a r t i c l e
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Article history: Received 25 June 2014 Accepted 31 December 2014 Available online 14 January 2015
a b s t r a c t A new phase was found to precipitate in a Fe–Cr model alloy after 58 keV deuterium ion irradiation at 773 K. The nanoscale radiation-induced precipitate was studied systematically using high resolution transmission electron microscopy (HRTEM), image simulation and in-situ ultrahigh voltage transmission electron microscopy (HVEM). B2 structure is proposed for the new Cr-rich phase, which adopts a cubeon-cube orientation relationship with regard to the Fe matrix. Geometric phase analysis (GPA) was employed to measure the strain fields around the precipitate and this was used to explain its characteristic 1-dimensional elongation along the h1 0 0i Fe direction. The precipitate was stable under subsequent electron irradiation at different temperatures. We suggest that the precipitate with a high interface-tovolume ratio enhances the radiation resistance of the material. The reason for this is the presence of a large number of interfaces between the precipitate and the matrix, which may greatly reduce the concentration of point defects around the dislocation loops. This leads to a significant decrease in the growth rate. Ó 2015 Elsevier B.V. All rights reserved.
1. Introduction Improved resistance to radiation damage is an important issue in the development of nuclear materials for fusion reactors as well as for the safe design and operation of innovative nuclear systems [1–5]. Nanostructured materials have been reported to have improved radiation resistance because of the presence of a large fraction of nano-interfaces [6–13]. An interface is defined as an area where two dissimilar chemical phases interact with each other [14]. From a traditional point of view, the interface can act as a sink for the annihilation of point defects (interstitials and vacancies) that are created during irradiation. Bai et al. recently reported that grain boundaries, which are prevalent in nanomaterials, can capture interstitials and transfer them back to the lattice leading to the annihilation of any vacancy that comes within a few ⇑ Corresponding author. Tel./fax: +86 010 62333580. E-mail address: [email protected] (Q. Zhan). http://dx.doi.org/10.1016/j.jnucmat.2014.12.124 0022-3115/Ó 2015 Elsevier B.V. All rights reserved.
nanometers of the grain boundary [15]. This new explanation mechanism is considered to be general and may be applied at interfaces between different materials. The previous experimental results and the interface-mechanism proposed suggest a new direction for improving radiation tolerance of material: increasing nano-interfaces. The presence of small-size precipitates in a material leads to a high interface-to-volume ratio. The inclusion of high-density nano size precipitates are considered to be effective for significant improvement of radiation resistance of the material. Effort has been put into increasing the amount of stabilized fine precipitate by controlling the alloying elements, heat-treatment or by adding a thermally stable oxide particle dispersion into the matrix. This is done to improve the performance of reduced activation ferritic/martensitic (RAFM) steels which are considered as candidate materials for advanced nuclear applications [8,16]. The reported precipitate can also be obtained by irradiation. Since the 1970s many radiation-induced precipitation (RIP) and segregation (RIS)
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phenomena have been observed in model alloys or in the structural materials of nuclear reactors [17,18]. RIS and RIP were first discovered by Okamoto and Wiedersich [19]. They observed the precipitation of an intermetallic compound at the surfaces of radiationinduced voids in an austenitic stainless steel. Shortly after this discovery Barbu and Ardell observed precipitation in a model alloy at interstitial dislocation loops as a consequence of ion irradiation at 773 K [20]. At atomic scale the origin of RIS and RIP is relatively well understood; they come from coupling between the fluxes of point defects created by irradiation and chemical species fluxes [17,21]. In the simplest possible terms, the fluxes of atoms that are induced by the fluxes of point defects lead to the radiationinduced segregation of the solute (and solvent) atoms and this is RIS. When the concentration of the solute exceeds its solubility limit a new phase will nucleate and grow leading to RIP. RIS has been widely used for decades to explain changes in elemental composition caused by irradiation. However, few studies exist on RIP in RAFM steel and the chemical components, the crystal structure, orientation relationships, stability under further irradiation and other effects on the irradiation behavior remain unclear. Recently, Arakawa reported that the segregation of chromium (Cr) occurs at the periphery of the loops in a Fe–Cr alloy under electron irradiation and at higher than 450 K [22]. A question thus arises about the fate of the irradiation induced nano-precipitate. In this paper, we present observations and a detailed characterization of a new Cr-rich precipitate that was induced by ion irradiation in a model alloy. Further study was motivated by questions about the stability of the precipitate under subsequent irradiation, and their effect on the irradiation performance of the materials. We discuss the dynamic growth behavior of defects that were preinduced by ion irradiation using in situ ultrahigh voltage transmission electron microscopy (HVTEM) data. The role of nano-precipitates is discussed in detail. 2. Experimental The material analyzed in this study was Fe–10 wt%Cr, which is a model alloy of reduced activation ferritic/martensitic (RAFM) steels [23–28]. Samples suitable for TEM studies were prepared using the standard techniques described elsewhere [29,30]. The TEM samples were mounted on a pure-Cu sample holder to ensure good thermal and electrical conductivity. This was followed by irradiation with deuterium ions in a LC-4 ion accelerator. Deuterium ions with an energy of 58 keV were implanted directly into the samples at room temperature (RT) and 773 K, respectively. 2 The ion fluence was 1.6 1021 Dþ 2 /m . The radiation damage induced by the ions [in units of displacements/atom (dpa)] was estimated by the Monte Carlo code SRIM 2008 [31]. The peak displacement damage dose was about 4.8 dpa based on the calculations (see Fig. 1). Microstructural investigations were carried out in a Tecnai F20 TEM. Electron-beam irradiation at different temperatures was carried out in an ultrahigh voltage TEM (H-3000) installed at Osaka University, Japan. The accelerator voltage of electron was 2 MV and the beam flux was 1.97 1023 e/m2 s. Irradiation damage rate was the product of the beam flux and the scattering cross section which is about 4.0 1027 (m2) for Fe–Cr sample according to the calculation [32,33]. Therefore the irradiation damage rate was 7.88 104 dpa/s. 3. Experimental results 3.1. Characterization of the new phase 3.1.1. Precipitation in the irradiated alloy TEM observations reveal that the materials have different microstructural features under irradiation at different
Fig. 1. Displacement damage profile and deuterium content profile in the sample calculated by SRIM. Calculation shown is for a fluence of 1.6 1021 ions/m2 at 58 keV.
temperatures. Fig. 2 shows low-magnification images of the Fe– 10Cr alloy with and without deuterium ion irradiation at RT and at 773 K. We found that the original sample was free from second phases as no additional diffraction spots to those of the matrix spots were observed, as shown in the inset electron diffraction pattern (EDP) (Fig. 2a). Conversely, many tiny black spots appeared in the samples upon RT deuterium-irradiation (Fig. 2b). No additional diffraction spots were present in the inset EDP and this confirms that the tiny black spots are not precipitates but irradiation defects such as dislocation loops [30]. A large amount of mutually perpendicular 2D ‘‘stripes’’ were present in the specimens that were irradiated with deuterium ions at 773 K, as shown in Fig. 2c. The corresponding EDP is given in the inset. Additional weak reflections in the middle of the main diffraction spots along the (2 0 0)Fe plane of the Fe matrix can be observed along the [0 0 1] zone axis. Fig. 2d shows corresponding dark field morphology images using superlattice diffraction spots. These confirm that the additional diffraction spots arise from the stripe precipitates. They extend along the h1 0 0i direction with a width of about 2– 8 nm and a length of tens to hundreds of nm. This remarkable feature shows that the high density nano-precipitate was indeed induced by deuterium ion irradiation at 773 K. 3.1.2. Structure of the new phase To investigate the structure of the precipitate a series of EDPs with different zone axes were obtained from the same area of the irradiated sample. These weak superlattice spots in the [0 0 1] direction were always present upon tilting. However, no existing phases could be indexed to the obtained stripe precipitate. These included FeCr, FexOy, CrxOy and Fe3C. We concluded that a new phase was produced because of the ion irradiation at 773 K. The crystal structure of the new phase was identified and determined using the recorded diffraction patterns. A series of EDPs 1 1] were obtained with different zone axes of [0 0 1], [0 1 1] and [1 from the irradiated sample, as shown in Fig. 3a–c. The intense reflections that belong to the Fe matrix were superimposed by weaker reflections from the precipitates. Both the Fe matrix and the precipitates show 4-fold and 3-fold axes in the [0 0 1] and [1 1 1] directions, respectively. This is consistent with their cubic symmetry. Considering the weak spots we propose a B2 structure for the new phase, which is an ordered bcc structure. From calculations carried out on these patterns the lattice parameter of the new phase embedded in the Fe matrix is exactly the same as that
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Fig. 2. Micro morphology of the Fe–10Cr model alloy without deuterium ion pre-irradiation (a). Bright field image of the sample with deuterium ion irradiation at RT (b). Bright field image (c) and dark field image (d) of the sample after ion irradiation at 773 K.
1 1] Fe Fig. 3. Electron diffraction patterns (EDPs) of the Fe–10Cr sample irradiated by deuterium at 773 K: (a) [0 0 1] Fe matrix projection, (b) [0 1 1] Fe projection, (c) [1 structure and with (d) [0 0 1], (e) [0 1 1], projection. In these EDPs the precipitate spots with the same orientation were verified by simulated diffraction patterns using a Pm3m 1 1] zone axes. and (f) [1
of Fe (ap = aFe, subscript p represents the precipitate), which is 0.2860 nm. The corresponding simulated EDPs for the cubic B2 space group and a lattice parameter of structure with the Pm3m 0.2860 nm are shown in Fig. 3d–f, respectively. The simulated results agree well with the experimental EDPs and this confirms our prediction of the crystal structure of the new phase. The origin of the other weak spots, apart from those with the B2 structure
shown in Fig. 3b, is still unclear. This is possibly due to a shape effect of the 2D stripe precipitate. The investigated EDPs are composed of two sets of patterns that belong to the new phase (weak reflections) and the Fe matrix (strong reflections), respectively (Fig. 3a–c). Cube-on-cube orientation relationships can thus be obtained between the Fe matrix and the new phase as follows: [0 0 1]pk[0 0 1]Fe, (1 1 0)pk(1 1 0)Fe.
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3.1.3. Composition of the precipitate The chromium content in the Fe matrix and the stripe precipitate were analyzed statistically using energy dispersive X-ray spectroscopy (EDX) at a high spatial resolution. The results are given in Fig. 4. Although a certain experimental error exists in quantitative EDX results because of the magnetic characteristic of Fe–Cr alloys, the distribution trend for chromium could still be obtained. The chromium content in the mutually perpendicular stripe precipitate was about twice that in the matrix and the precipitate was thus rich in chromium. Classical RIP theories [17,34] may be used to explain the phase precipitation observed in this study. The atomic radius of Cr is slightly larger than that of Fe. Molecular dynamics (MD) simulation work reveals that complex Cr behavior exists in bcc Fe–Cr systems [35]. The nearest neighbors may relax towards the Cr atoms resulting in a change in Cr from oversized to undersized. This depends on the concentration and electronic/magnetic effects. Therefore, for undersized Cr atoms, Cr-self-interstitial Fe complexes may form preferentially because of the resultant positive binding energy. Cr atoms will thus be solute-dragged towards the dislocation loops. When the Cr atom concentration slightly exceeds its solubility
limit in Fe, a new Cr-rich phase will nucleate and grow. It is worth noting that strong interactions among Cr atoms, self-interstitial atoms [36,37] and deuterium implanted at 773 K probably play an important role in the appearance of the new phase. To summarize the above-mentioned observations, a new Crrich phase is induced by deuterium ion irradiation and it homogeneously precipitates in the Fe–Cr model alloy. The high density nano-precipitate has a cubic structure and a cube-on-cube orientation relationship with the Fe matrix. 3.2. Measurement of strain fields around the precipitate A typical high resolution TEM (HRTEM) image of the stripe precipitate is given in Fig. 5. A stripe of 2–5 nm in width and up to about 50 nm in length is clearly present. The incident beam was along the [0 0 1] zone axis. A fast Fourier transformation (FFT) of the HRTEM image is shown in Fig. 5b. The superlattice diffraction spots and the streak diffraction are distinct. The strain fields around the precipitate can be measured and mapped from the HRTEM image of the precipitated phase using the geometric phase analysis (GPA) technique [38–40]. Fig. 5c shows a local magnified
Fig. 4. Statistics of the Cr content in matrix (a) and precipitate (b) in the Fe–10Cr irradiated sample.
Fig. 5. High resolution electron microscope image of a stripe precipitate. (a) HRTEM image of the precipitate, (b) FFT image of the HRTEM image, (c) enlarged square region shows lattice fringes and (d) FFT image of (c).
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Fig. 6. Geometric phase analysis of the observed stripe precipitate in the model alloy. Strain maps (above) and profiles (below) of the precipitate are shown here.
image of this precipitate. Note that the image was recorded on using a CCD with 2048 by 2048 pixel CCD and the numerical images have a low sampling density of 0.015 nm/pixel. This which corresponds to only about 10 pixels per (200)Fe lattice fringe. Therefore, it is possible to accurately measure and map the strain fields around the precipitate by GPA. We selected the two intense spots in the FFT image (Fig. 5d) in Fig. 5c at, g1 = (1 0 0)P, which arose from the precipitate, and at g2 = (1 1 0)Fe as marked to calculate the strain field. Because the intensity of the diffraction spots from the precipitate was not very strong the results may not express the best signal-to-noise ratio. We defined the x-axis as being parallel to the [0 2 0] direction of the Fe matrix and the y-axis as the [2 0 0] direction. The normal strain along the x-axis, exx, and along the y-axis, eyy, and the shear strain exy are shown in Fig. 6a–c,
respectively. We used identical color scales. Around the precipitate we obtained the following average strain values: exx of 0.45 ± 0.2%, exy of 0.2 ± 0.2% and eyy along the elongation direction of the stripe precipitate of 0.01 ± 0.2%, and this was the smallest value. These values explain why the new phase grew along the h1 0 0i direction. The profiles across the precipitate were extracted to show the quantitative value of the strain field, as shown in Fig. 6d–f. An obvious change in the strain across the interface between the precipitate and the matrix was observed. 3.3. Stability of the precipitate under electron irradiation An in situ electron irradiation experiment was then performed on Fe–10Cr specimens that were irradiated with deuterium ions.
Fig. 7. In situ observation of the Cr-rich precipitate in the model alloy upon deuterium ion irradiation at 773 K and under electron irradiation at RT.
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Fig. 8. In situ observation of the Cr-rich phase in the model alloy pre-irradiated with deuterium ions under electron irradiation at 773 K.
This was done to investigate the stability of the Cr-rich stripe precipitate. The electron irradiation temperatures were RT and 773 K. Fig. 7 contains a series of sequential micrographs that show the microstructural change in the model samples under electron irradiation at RT. The original lengths of the precipitate, marked with white line segments, were about 96.1 nm and 155.2 nm, respectively. The original size (length of the long axis) of the dislocation loop indicated by the black arrow is about 29.9 nm. Under continuous electron irradiation an increasing number of new defects and defect clusters emerged and the dislocation loop grew significantly. With electron irradiation of 0.17 dpa the size of the dislocation loop was about 42.6 nm. However, the length of the marked precipitates did not change obviously with a dose of up to 0.58 dpa. Further irradiation led to the dislocation loop becoming slowly blurred because the defects or defect clusters became tangled. No distinct change was observed for the precipitates. The electron irradiation was also carried out in a similar manner at 773 K, as shown in Fig. 8. No significant change in size was observed for the Cr-rich precipitates induced by high temperature deuterium ion irradiation. This indicates their high stability under electron irradiation, even at a high temperature of 773 K. The in situ electron diffraction pattern of the same irradiation area is also given in Fig. 8 after an irradiation dose of 2.13 dpa. The
diffraction spots of the precipitates are clearly still present suggesting that the crystallographic structure remained unchanged. The dislocation loops grew significantly upon electron irradiation. The dependence of precipitate size on irradiation is shown in Fig. 9. A slight change in precipitate size is evident within several nanometers. From in situ observations, the Cr-rich precipitate retains a stable size and structure under the subsequent electron irradiation. 3.4. Effects of the precipitate on dislocation loop growth The in situ electron irradiation experiment was also used to investigate the effects of the nano-precipitate on the growth of the dislocation loop. The data was also used to investigate the dynamic behavior of the point defects [interstitial atoms (I) and vacancy (V)] during electron irradiation. The physical process is discussed in detail as follows. The dislocation loops grew in a distinct manner upon electron irradiation either at RT or at 773 K. We know that the dislocation loop will grow faster at higher temperatures [30]. To better compare and clarify the experimental results, Fig. 10 shows the dislocation loop growth process under electron irradiation at 773 K in the specimens irradiated by deuterium ions at RT where no precipitate was induced. Note that the same electron irradiation
Fig. 9. Irradiation dependence of precipitates size: (a) electron irradiation at RT and (b) electron irradiation at 773 K.
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conditions were used for both types of specimens with or without the induction of precipitates. Dislocation loops with a similar original size were selected for the following analysis. The growth of the loops under electron irradiation at 773 K in the samples with and without irradiation induced precipitation is shown in Fig. 11. Both the dislocation loops in the different samples grew with irradiation but their growth rates were quite different. In the sample without the pre-irradiation induced precipitate, the growth rate of the dislocation loop was about 125.6 nm/dpa at a radiation dose of 0.17 dpa. We obtained just 5.5 nm/dpa in the sample containing precipitates with the same electron irradiation dose. At an irradiation dose of 2.13 dpa, the growth rate of the dislocation loop with the precipitate was 5.4 nm/dpa. We thus suggest that the dislocation loop will grow far slower because of the presence of nano-precipitate with high density. To further exclude the effect of the original size of the dislocation loops on the growth rate we defined the growth ratio of the dislocation loop as:
s ¼ ð½d d0 =d0 Þ 100%
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Fig. 11. Dislocation loop size as a function of irradiation dose.
ð1Þ
where d0 is the original size of the dislocation loop and d is the size of the loop under irradiation. In the sample without a precipitate the growth ratio of the dislocation loop calculated from Fig. 11 was 98.9% upon 0.17 dpa irradiation at 773 K. A growth ratio of only 3.9% was measured in the sample with precipitates under the same irradiation dose and temperature. This ratio increased to 49.6% with about 2.13 dpa irradiation at the same temperature. The results reveal that the growth ratio of the dislocation loops can be greatly reduced owing to the presence of nanoscale precipitates. 4. Discussion High voltage electron irradiation produces V and I in equal numbers and their concentrations far exceed their equilibrium values. Healing occurs when they mostly recombine and annihilate. The rest single point defects attract each other and form an interstitial atom cluster (Is) and a vacancy cluster (Vs). The movement and aggregation of point defects in the general crystalline material is shown in Fig. 12a. I and Is complexes may move preferentially while leaving vacancies behind because interstitial atoms migrate
more easily than vacancies [10,36]. The complexes are usually trapped at the surface and at the interstitial loop sites [41–43] that were induced by ion pre-irradiation. This causes the general crystalline material to swell leading to the fast growth of loops as observed in our experiment. When the irradiation temperature is increased, V begins to migrate freely and so do the Vs complexes. Voids may form when V or related Vs combine and grow. The dislocation loops cause irradiation hardening and the voids may cause the materials to swell. When a secondary phase precipitates in the matrix of the crystalline material, an interface forms and acts as a sink. This affects the movement and aggregation of the point defects (Fig. 12b). As the number of sinks increase, the I and Is complexes are less likely to be trapped at the dislocation loops and this affects the growth of the loops. This is also the case for the V and Vs complexes. Therefore we investigated the growth process of the loops and voids to understand the influence of the Cr-rich precipitate on the movement of interstitial atoms and vacancies and to assess variations in the irradiation property. On the whole, irradiation-induced processes, though far from equilibrium, can be treated as a set of large number of local
Fig. 10. Dislocation loop growth process in the Fe–Cr model alloy irradiated with deuterium ions at RT under electrons irradiation at 773 K.
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Fig. 12. Schematic of the movement and aggregation of point defects. In a general crystalline material (a), the interstitial defects caused by irradiation quickly move to the surface and dislocation loops causing growth of loop, excepting most of defects annihilated. Vacancies slowly agglomerate and form immobile voids that cause the material to become brittle and to swell. In a material with a precipitate in matrix (b), the point defects get stuck at the interface. This greatly reduces the defect concentration in the matrix, which enhances the irradiation resistance of the materials.
quasi-equilibrium state if each volume is small enough. In order to understand the effect of precipitate on the interstitial atoms and surrounding dislocation loops, an assumption has been proposed that a single precipitate, dislocation loop and surrounding several nm regions is a quasi-equilibrium system. And the whole nonequilibrium material is composed of large number of such regional quasi-equilibrium regions. Based on the assumptions, the following thermodynamic equation was used to understand the effects of the precipitate qualitatively. The relationship between the concentration of the interstitial atoms CI and the interaction energy of the sink and the interstitial atoms can be described as follows (similar situations for vacancies are not shown here):
U L I C LI ¼ C 0 exp kT
ð3Þ
ð4Þ
From Eqs. (3) and (4) the interstitial concentrations can be calculated as follows:
0 U IFI U LI C LI ¼ C 0 1 exp exp kT kT
where R is the radius of the dislocation loop, ZIL and ZVL are absorption cross-sections of the loops with point defects I and V, respectively, MI is the interstitial atom jump rate and b is a parameter related to the absorption cross-section. b is a constant (at least at a fixed temperature) and is thought to be slightly larger than unity. This is because of the larger capture cross-section ZIL for interstitials than that for vacancies ZVL around a dislocation. At a fixed temperature, MI is reasonably assumed to be constant. Therefore, the growth rate of the loops may be expressed as follows using the constant A to represent the coefficients in Eq. (6):
dR ¼ AC I dt
where C LI and C IF I indicate the concentrations of the interstitial atoms around the dislocation loop and the interface, respectively. C0 is the average defect concentration in the crystal grain. ULI and UIFI represent the interaction energy between the interstitial atoms and the loop and interface, respectively. k is Boltzmann’s constant and T is the thermodynamic temperature (K). In the general crystalline material the concentration of interstitial atoms around the dislocation loop can be described by Eq. (2). When a secondary phase precipitates in the matrix the concentration will decrease and can be described as:
0 U LI C LI ¼ C 0 C IF exp I kT
ð6Þ
ð2Þ
and
U IFI C IF I ¼ C 0 exp kT
dR ¼ aðZ IL bZ VL ÞMI C I dt
ð5Þ
This means that the interstitial concentration will decrease significantly if a large number of interfaces, which may be introduced by the nano-precipitates, form in the matrix of the materials. The rate of growth of the loop can be expressed as [44]:
ð7Þ
Here, the growth rate of the interstitial loop is controlled by the concentration C LI of the interstitial atoms around the dislocation loop. The interface between the precipitate and the matrix is an important factor that affects the concentration and the movement of point defects. This may provide vacancy-interstitial recombination centers that heal defects in the material. Therefore, the presence of high density precipitates will lead to a decrease in the growth rate of the dislocation loops by reducing the concentration of interstitial atoms around the dislocation loops and this may decrease the production of voids in these materials because of the annihilation of vacancy-interstitials at the interface. An analysis of the above-mentioned results reveal that the growth rate of dislocation loops was about 125.6 nm/dpa in the sample without the precipitate under 0.17 dpa irradiation. The growth rate was just 5.5 nm/dpa in the sample with precipitates. This means that the induced precipitates in this experiment greatly reduce the concentration of interstitial atoms around the dislocation loop. This leads to a significant decrease in the growth rate. In our experiment, no distinct void was found, even at a high temperature of 773 K with an irradiation dose of up to 2.13 dpa. Therefore, the nanoscale precipitates with high density that were induced by suitable pre-ion irradiation may lead to an improvement in the irradiation performance of these materials. This work thus provides a simple method to explore anti-irradiation in RAFM materials.
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5. Conclusions In addition to dislocation loops, irradiation with deuterium ions at 773 K induced the precipitation of a new phase rich in Cr in the Fe–10Cr model alloy. These nanoscale precipitates were arranged along the h1 0 0i direction with widths of 2–8 nm and lengths ranging from tens to hundreds of nm. The 1-D elongation of the precipitate is attributed to the distribution of strain fields around the precipitates. The results from a series of EDPs combined with diffraction calculations reveal that the phase crystallizes in the cubic space group and adopts a cube-onsystem. It belongs to the Pm3m cube orientation relationship with the Fe matrix and it has the same lattice parameters as Fe. Under subsequent 2 MV electron irradiation the precipitates are highly stable while the surrounding dislocation loops grow far more slowly compared with the loops in the specimen without ion irradiation induced precipitates. We argue that the distinct decrease in the growth rate of dislocation loops comes from the presence of a large number of fine precipitates. This is because the interfaces between the precipitate and the matrix may provide vacancy-interstitial recombination centers to heal defects in the material. We suggest that this precipitation at the nanoscale that is induced by deuterium ion irradiation may enhance the irradiation resistance of a material because of a reduction in the concentration of point defects. Acknowledgements This work is supported by the National Magnetic Confinement Fusion Program (Grant Nos. 2011GB108002, 2014GB104003 and 2014GB120001), the National Natural Science Foundation of China (Grant No. 51371031). The authors gratefully acknowledge the Research Center for Ultra-High Voltage Electron Microscopy at Osaka University, Japan. The help and suggestions from Professor H. Mori and Dr. K. Arakawa at Osaka University were highly appreciated. References [1] [2] [3] [4] [5]
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Deuterium ion irradiation induced precipitation in Fe–Cr alloy: Characterization and effects on irradiation behavior P.P. Liu a, R. Yu b, Y.M. Zhu a, M.Z. Zhao a, J.W. Bai a, F.R. Wan a, Q. Zhan a,⇑ a b
School of Materials Science and Engineering, University of Science and Technology Beijing, Beijing 100083, China Department of Materials Science and Engineering, Beijing National Center for Electron Microscopy, Tsinghua University, Beijing 100084, China
h i g h l i g h t s A new phase precipitated in Fe–Cr alloy after deuterium ion irradiation at 773 K. B2 structure was proposed for the Cr-rich new phase. Strain fields around the precipitate have been measured by GPA. The precipitate decrease growth rate of dislocation loop under electron irradiation.
a r t i c l e
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Article history: Received 25 June 2014 Accepted 31 December 2014 Available online 14 January 2015
a b s t r a c t A new phase was found to precipitate in a Fe–Cr model alloy after 58 keV deuterium ion irradiation at 773 K. The nanoscale radiation-induced precipitate was studied systematically using high resolution transmission electron microscopy (HRTEM), image simulation and in-situ ultrahigh voltage transmission electron microscopy (HVEM). B2 structure is proposed for the new Cr-rich phase, which adopts a cubeon-cube orientation relationship with regard to the Fe matrix. Geometric phase analysis (GPA) was employed to measure the strain fields around the precipitate and this was used to explain its characteristic 1-dimensional elongation along the h1 0 0i Fe direction. The precipitate was stable under subsequent electron irradiation at different temperatures. We suggest that the precipitate with a high interface-tovolume ratio enhances the radiation resistance of the material. The reason for this is the presence of a large number of interfaces between the precipitate and the matrix, which may greatly reduce the concentration of point defects around the dislocation loops. This leads to a significant decrease in the growth rate. Ó 2015 Elsevier B.V. All rights reserved.
1. Introduction Improved resistance to radiation damage is an important issue in the development of nuclear materials for fusion reactors as well as for the safe design and operation of innovative nuclear systems [1–5]. Nanostructured materials have been reported to have improved radiation resistance because of the presence of a large fraction of nano-interfaces [6–13]. An interface is defined as an area where two dissimilar chemical phases interact with each other [14]. From a traditional point of view, the interface can act as a sink for the annihilation of point defects (interstitials and vacancies) that are created during irradiation. Bai et al. recently reported that grain boundaries, which are prevalent in nanomaterials, can capture interstitials and transfer them back to the lattice leading to the annihilation of any vacancy that comes within a few ⇑ Corresponding author. Tel./fax: +86 010 62333580. E-mail address: [email protected] (Q. Zhan). http://dx.doi.org/10.1016/j.jnucmat.2014.12.124 0022-3115/Ó 2015 Elsevier B.V. All rights reserved.
nanometers of the grain boundary [15]. This new explanation mechanism is considered to be general and may be applied at interfaces between different materials. The previous experimental results and the interface-mechanism proposed suggest a new direction for improving radiation tolerance of material: increasing nano-interfaces. The presence of small-size precipitates in a material leads to a high interface-to-volume ratio. The inclusion of high-density nano size precipitates are considered to be effective for significant improvement of radiation resistance of the material. Effort has been put into increasing the amount of stabilized fine precipitate by controlling the alloying elements, heat-treatment or by adding a thermally stable oxide particle dispersion into the matrix. This is done to improve the performance of reduced activation ferritic/martensitic (RAFM) steels which are considered as candidate materials for advanced nuclear applications [8,16]. The reported precipitate can also be obtained by irradiation. Since the 1970s many radiation-induced precipitation (RIP) and segregation (RIS)
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phenomena have been observed in model alloys or in the structural materials of nuclear reactors [17,18]. RIS and RIP were first discovered by Okamoto and Wiedersich [19]. They observed the precipitation of an intermetallic compound at the surfaces of radiationinduced voids in an austenitic stainless steel. Shortly after this discovery Barbu and Ardell observed precipitation in a model alloy at interstitial dislocation loops as a consequence of ion irradiation at 773 K [20]. At atomic scale the origin of RIS and RIP is relatively well understood; they come from coupling between the fluxes of point defects created by irradiation and chemical species fluxes [17,21]. In the simplest possible terms, the fluxes of atoms that are induced by the fluxes of point defects lead to the radiationinduced segregation of the solute (and solvent) atoms and this is RIS. When the concentration of the solute exceeds its solubility limit a new phase will nucleate and grow leading to RIP. RIS has been widely used for decades to explain changes in elemental composition caused by irradiation. However, few studies exist on RIP in RAFM steel and the chemical components, the crystal structure, orientation relationships, stability under further irradiation and other effects on the irradiation behavior remain unclear. Recently, Arakawa reported that the segregation of chromium (Cr) occurs at the periphery of the loops in a Fe–Cr alloy under electron irradiation and at higher than 450 K [22]. A question thus arises about the fate of the irradiation induced nano-precipitate. In this paper, we present observations and a detailed characterization of a new Cr-rich precipitate that was induced by ion irradiation in a model alloy. Further study was motivated by questions about the stability of the precipitate under subsequent irradiation, and their effect on the irradiation performance of the materials. We discuss the dynamic growth behavior of defects that were preinduced by ion irradiation using in situ ultrahigh voltage transmission electron microscopy (HVTEM) data. The role of nano-precipitates is discussed in detail. 2. Experimental The material analyzed in this study was Fe–10 wt%Cr, which is a model alloy of reduced activation ferritic/martensitic (RAFM) steels [23–28]. Samples suitable for TEM studies were prepared using the standard techniques described elsewhere [29,30]. The TEM samples were mounted on a pure-Cu sample holder to ensure good thermal and electrical conductivity. This was followed by irradiation with deuterium ions in a LC-4 ion accelerator. Deuterium ions with an energy of 58 keV were implanted directly into the samples at room temperature (RT) and 773 K, respectively. 2 The ion fluence was 1.6 1021 Dþ 2 /m . The radiation damage induced by the ions [in units of displacements/atom (dpa)] was estimated by the Monte Carlo code SRIM 2008 [31]. The peak displacement damage dose was about 4.8 dpa based on the calculations (see Fig. 1). Microstructural investigations were carried out in a Tecnai F20 TEM. Electron-beam irradiation at different temperatures was carried out in an ultrahigh voltage TEM (H-3000) installed at Osaka University, Japan. The accelerator voltage of electron was 2 MV and the beam flux was 1.97 1023 e/m2 s. Irradiation damage rate was the product of the beam flux and the scattering cross section which is about 4.0 1027 (m2) for Fe–Cr sample according to the calculation [32,33]. Therefore the irradiation damage rate was 7.88 104 dpa/s. 3. Experimental results 3.1. Characterization of the new phase 3.1.1. Precipitation in the irradiated alloy TEM observations reveal that the materials have different microstructural features under irradiation at different
Fig. 1. Displacement damage profile and deuterium content profile in the sample calculated by SRIM. Calculation shown is for a fluence of 1.6 1021 ions/m2 at 58 keV.
temperatures. Fig. 2 shows low-magnification images of the Fe– 10Cr alloy with and without deuterium ion irradiation at RT and at 773 K. We found that the original sample was free from second phases as no additional diffraction spots to those of the matrix spots were observed, as shown in the inset electron diffraction pattern (EDP) (Fig. 2a). Conversely, many tiny black spots appeared in the samples upon RT deuterium-irradiation (Fig. 2b). No additional diffraction spots were present in the inset EDP and this confirms that the tiny black spots are not precipitates but irradiation defects such as dislocation loops [30]. A large amount of mutually perpendicular 2D ‘‘stripes’’ were present in the specimens that were irradiated with deuterium ions at 773 K, as shown in Fig. 2c. The corresponding EDP is given in the inset. Additional weak reflections in the middle of the main diffraction spots along the (2 0 0)Fe plane of the Fe matrix can be observed along the [0 0 1] zone axis. Fig. 2d shows corresponding dark field morphology images using superlattice diffraction spots. These confirm that the additional diffraction spots arise from the stripe precipitates. They extend along the h1 0 0i direction with a width of about 2– 8 nm and a length of tens to hundreds of nm. This remarkable feature shows that the high density nano-precipitate was indeed induced by deuterium ion irradiation at 773 K. 3.1.2. Structure of the new phase To investigate the structure of the precipitate a series of EDPs with different zone axes were obtained from the same area of the irradiated sample. These weak superlattice spots in the [0 0 1] direction were always present upon tilting. However, no existing phases could be indexed to the obtained stripe precipitate. These included FeCr, FexOy, CrxOy and Fe3C. We concluded that a new phase was produced because of the ion irradiation at 773 K. The crystal structure of the new phase was identified and determined using the recorded diffraction patterns. A series of EDPs 1 1] were obtained with different zone axes of [0 0 1], [0 1 1] and [1 from the irradiated sample, as shown in Fig. 3a–c. The intense reflections that belong to the Fe matrix were superimposed by weaker reflections from the precipitates. Both the Fe matrix and the precipitates show 4-fold and 3-fold axes in the [0 0 1] and [1 1 1] directions, respectively. This is consistent with their cubic symmetry. Considering the weak spots we propose a B2 structure for the new phase, which is an ordered bcc structure. From calculations carried out on these patterns the lattice parameter of the new phase embedded in the Fe matrix is exactly the same as that
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Fig. 2. Micro morphology of the Fe–10Cr model alloy without deuterium ion pre-irradiation (a). Bright field image of the sample with deuterium ion irradiation at RT (b). Bright field image (c) and dark field image (d) of the sample after ion irradiation at 773 K.
1 1] Fe Fig. 3. Electron diffraction patterns (EDPs) of the Fe–10Cr sample irradiated by deuterium at 773 K: (a) [0 0 1] Fe matrix projection, (b) [0 1 1] Fe projection, (c) [1 structure and with (d) [0 0 1], (e) [0 1 1], projection. In these EDPs the precipitate spots with the same orientation were verified by simulated diffraction patterns using a Pm3m 1 1] zone axes. and (f) [1
of Fe (ap = aFe, subscript p represents the precipitate), which is 0.2860 nm. The corresponding simulated EDPs for the cubic B2 space group and a lattice parameter of structure with the Pm3m 0.2860 nm are shown in Fig. 3d–f, respectively. The simulated results agree well with the experimental EDPs and this confirms our prediction of the crystal structure of the new phase. The origin of the other weak spots, apart from those with the B2 structure
shown in Fig. 3b, is still unclear. This is possibly due to a shape effect of the 2D stripe precipitate. The investigated EDPs are composed of two sets of patterns that belong to the new phase (weak reflections) and the Fe matrix (strong reflections), respectively (Fig. 3a–c). Cube-on-cube orientation relationships can thus be obtained between the Fe matrix and the new phase as follows: [0 0 1]pk[0 0 1]Fe, (1 1 0)pk(1 1 0)Fe.
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3.1.3. Composition of the precipitate The chromium content in the Fe matrix and the stripe precipitate were analyzed statistically using energy dispersive X-ray spectroscopy (EDX) at a high spatial resolution. The results are given in Fig. 4. Although a certain experimental error exists in quantitative EDX results because of the magnetic characteristic of Fe–Cr alloys, the distribution trend for chromium could still be obtained. The chromium content in the mutually perpendicular stripe precipitate was about twice that in the matrix and the precipitate was thus rich in chromium. Classical RIP theories [17,34] may be used to explain the phase precipitation observed in this study. The atomic radius of Cr is slightly larger than that of Fe. Molecular dynamics (MD) simulation work reveals that complex Cr behavior exists in bcc Fe–Cr systems [35]. The nearest neighbors may relax towards the Cr atoms resulting in a change in Cr from oversized to undersized. This depends on the concentration and electronic/magnetic effects. Therefore, for undersized Cr atoms, Cr-self-interstitial Fe complexes may form preferentially because of the resultant positive binding energy. Cr atoms will thus be solute-dragged towards the dislocation loops. When the Cr atom concentration slightly exceeds its solubility
limit in Fe, a new Cr-rich phase will nucleate and grow. It is worth noting that strong interactions among Cr atoms, self-interstitial atoms [36,37] and deuterium implanted at 773 K probably play an important role in the appearance of the new phase. To summarize the above-mentioned observations, a new Crrich phase is induced by deuterium ion irradiation and it homogeneously precipitates in the Fe–Cr model alloy. The high density nano-precipitate has a cubic structure and a cube-on-cube orientation relationship with the Fe matrix. 3.2. Measurement of strain fields around the precipitate A typical high resolution TEM (HRTEM) image of the stripe precipitate is given in Fig. 5. A stripe of 2–5 nm in width and up to about 50 nm in length is clearly present. The incident beam was along the [0 0 1] zone axis. A fast Fourier transformation (FFT) of the HRTEM image is shown in Fig. 5b. The superlattice diffraction spots and the streak diffraction are distinct. The strain fields around the precipitate can be measured and mapped from the HRTEM image of the precipitated phase using the geometric phase analysis (GPA) technique [38–40]. Fig. 5c shows a local magnified
Fig. 4. Statistics of the Cr content in matrix (a) and precipitate (b) in the Fe–10Cr irradiated sample.
Fig. 5. High resolution electron microscope image of a stripe precipitate. (a) HRTEM image of the precipitate, (b) FFT image of the HRTEM image, (c) enlarged square region shows lattice fringes and (d) FFT image of (c).
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Fig. 6. Geometric phase analysis of the observed stripe precipitate in the model alloy. Strain maps (above) and profiles (below) of the precipitate are shown here.
image of this precipitate. Note that the image was recorded on using a CCD with 2048 by 2048 pixel CCD and the numerical images have a low sampling density of 0.015 nm/pixel. This which corresponds to only about 10 pixels per (200)Fe lattice fringe. Therefore, it is possible to accurately measure and map the strain fields around the precipitate by GPA. We selected the two intense spots in the FFT image (Fig. 5d) in Fig. 5c at, g1 = (1 0 0)P, which arose from the precipitate, and at g2 = (1 1 0)Fe as marked to calculate the strain field. Because the intensity of the diffraction spots from the precipitate was not very strong the results may not express the best signal-to-noise ratio. We defined the x-axis as being parallel to the [0 2 0] direction of the Fe matrix and the y-axis as the [2 0 0] direction. The normal strain along the x-axis, exx, and along the y-axis, eyy, and the shear strain exy are shown in Fig. 6a–c,
respectively. We used identical color scales. Around the precipitate we obtained the following average strain values: exx of 0.45 ± 0.2%, exy of 0.2 ± 0.2% and eyy along the elongation direction of the stripe precipitate of 0.01 ± 0.2%, and this was the smallest value. These values explain why the new phase grew along the h1 0 0i direction. The profiles across the precipitate were extracted to show the quantitative value of the strain field, as shown in Fig. 6d–f. An obvious change in the strain across the interface between the precipitate and the matrix was observed. 3.3. Stability of the precipitate under electron irradiation An in situ electron irradiation experiment was then performed on Fe–10Cr specimens that were irradiated with deuterium ions.
Fig. 7. In situ observation of the Cr-rich precipitate in the model alloy upon deuterium ion irradiation at 773 K and under electron irradiation at RT.
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Fig. 8. In situ observation of the Cr-rich phase in the model alloy pre-irradiated with deuterium ions under electron irradiation at 773 K.
This was done to investigate the stability of the Cr-rich stripe precipitate. The electron irradiation temperatures were RT and 773 K. Fig. 7 contains a series of sequential micrographs that show the microstructural change in the model samples under electron irradiation at RT. The original lengths of the precipitate, marked with white line segments, were about 96.1 nm and 155.2 nm, respectively. The original size (length of the long axis) of the dislocation loop indicated by the black arrow is about 29.9 nm. Under continuous electron irradiation an increasing number of new defects and defect clusters emerged and the dislocation loop grew significantly. With electron irradiation of 0.17 dpa the size of the dislocation loop was about 42.6 nm. However, the length of the marked precipitates did not change obviously with a dose of up to 0.58 dpa. Further irradiation led to the dislocation loop becoming slowly blurred because the defects or defect clusters became tangled. No distinct change was observed for the precipitates. The electron irradiation was also carried out in a similar manner at 773 K, as shown in Fig. 8. No significant change in size was observed for the Cr-rich precipitates induced by high temperature deuterium ion irradiation. This indicates their high stability under electron irradiation, even at a high temperature of 773 K. The in situ electron diffraction pattern of the same irradiation area is also given in Fig. 8 after an irradiation dose of 2.13 dpa. The
diffraction spots of the precipitates are clearly still present suggesting that the crystallographic structure remained unchanged. The dislocation loops grew significantly upon electron irradiation. The dependence of precipitate size on irradiation is shown in Fig. 9. A slight change in precipitate size is evident within several nanometers. From in situ observations, the Cr-rich precipitate retains a stable size and structure under the subsequent electron irradiation. 3.4. Effects of the precipitate on dislocation loop growth The in situ electron irradiation experiment was also used to investigate the effects of the nano-precipitate on the growth of the dislocation loop. The data was also used to investigate the dynamic behavior of the point defects [interstitial atoms (I) and vacancy (V)] during electron irradiation. The physical process is discussed in detail as follows. The dislocation loops grew in a distinct manner upon electron irradiation either at RT or at 773 K. We know that the dislocation loop will grow faster at higher temperatures [30]. To better compare and clarify the experimental results, Fig. 10 shows the dislocation loop growth process under electron irradiation at 773 K in the specimens irradiated by deuterium ions at RT where no precipitate was induced. Note that the same electron irradiation
Fig. 9. Irradiation dependence of precipitates size: (a) electron irradiation at RT and (b) electron irradiation at 773 K.
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conditions were used for both types of specimens with or without the induction of precipitates. Dislocation loops with a similar original size were selected for the following analysis. The growth of the loops under electron irradiation at 773 K in the samples with and without irradiation induced precipitation is shown in Fig. 11. Both the dislocation loops in the different samples grew with irradiation but their growth rates were quite different. In the sample without the pre-irradiation induced precipitate, the growth rate of the dislocation loop was about 125.6 nm/dpa at a radiation dose of 0.17 dpa. We obtained just 5.5 nm/dpa in the sample containing precipitates with the same electron irradiation dose. At an irradiation dose of 2.13 dpa, the growth rate of the dislocation loop with the precipitate was 5.4 nm/dpa. We thus suggest that the dislocation loop will grow far slower because of the presence of nano-precipitate with high density. To further exclude the effect of the original size of the dislocation loops on the growth rate we defined the growth ratio of the dislocation loop as:
s ¼ ð½d d0 =d0 Þ 100%
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Fig. 11. Dislocation loop size as a function of irradiation dose.
ð1Þ
where d0 is the original size of the dislocation loop and d is the size of the loop under irradiation. In the sample without a precipitate the growth ratio of the dislocation loop calculated from Fig. 11 was 98.9% upon 0.17 dpa irradiation at 773 K. A growth ratio of only 3.9% was measured in the sample with precipitates under the same irradiation dose and temperature. This ratio increased to 49.6% with about 2.13 dpa irradiation at the same temperature. The results reveal that the growth ratio of the dislocation loops can be greatly reduced owing to the presence of nanoscale precipitates. 4. Discussion High voltage electron irradiation produces V and I in equal numbers and their concentrations far exceed their equilibrium values. Healing occurs when they mostly recombine and annihilate. The rest single point defects attract each other and form an interstitial atom cluster (Is) and a vacancy cluster (Vs). The movement and aggregation of point defects in the general crystalline material is shown in Fig. 12a. I and Is complexes may move preferentially while leaving vacancies behind because interstitial atoms migrate
more easily than vacancies [10,36]. The complexes are usually trapped at the surface and at the interstitial loop sites [41–43] that were induced by ion pre-irradiation. This causes the general crystalline material to swell leading to the fast growth of loops as observed in our experiment. When the irradiation temperature is increased, V begins to migrate freely and so do the Vs complexes. Voids may form when V or related Vs combine and grow. The dislocation loops cause irradiation hardening and the voids may cause the materials to swell. When a secondary phase precipitates in the matrix of the crystalline material, an interface forms and acts as a sink. This affects the movement and aggregation of the point defects (Fig. 12b). As the number of sinks increase, the I and Is complexes are less likely to be trapped at the dislocation loops and this affects the growth of the loops. This is also the case for the V and Vs complexes. Therefore we investigated the growth process of the loops and voids to understand the influence of the Cr-rich precipitate on the movement of interstitial atoms and vacancies and to assess variations in the irradiation property. On the whole, irradiation-induced processes, though far from equilibrium, can be treated as a set of large number of local
Fig. 10. Dislocation loop growth process in the Fe–Cr model alloy irradiated with deuterium ions at RT under electrons irradiation at 773 K.
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Fig. 12. Schematic of the movement and aggregation of point defects. In a general crystalline material (a), the interstitial defects caused by irradiation quickly move to the surface and dislocation loops causing growth of loop, excepting most of defects annihilated. Vacancies slowly agglomerate and form immobile voids that cause the material to become brittle and to swell. In a material with a precipitate in matrix (b), the point defects get stuck at the interface. This greatly reduces the defect concentration in the matrix, which enhances the irradiation resistance of the materials.
quasi-equilibrium state if each volume is small enough. In order to understand the effect of precipitate on the interstitial atoms and surrounding dislocation loops, an assumption has been proposed that a single precipitate, dislocation loop and surrounding several nm regions is a quasi-equilibrium system. And the whole nonequilibrium material is composed of large number of such regional quasi-equilibrium regions. Based on the assumptions, the following thermodynamic equation was used to understand the effects of the precipitate qualitatively. The relationship between the concentration of the interstitial atoms CI and the interaction energy of the sink and the interstitial atoms can be described as follows (similar situations for vacancies are not shown here):
U L I C LI ¼ C 0 exp kT
ð3Þ
ð4Þ
From Eqs. (3) and (4) the interstitial concentrations can be calculated as follows:
0 U IFI U LI C LI ¼ C 0 1 exp exp kT kT
where R is the radius of the dislocation loop, ZIL and ZVL are absorption cross-sections of the loops with point defects I and V, respectively, MI is the interstitial atom jump rate and b is a parameter related to the absorption cross-section. b is a constant (at least at a fixed temperature) and is thought to be slightly larger than unity. This is because of the larger capture cross-section ZIL for interstitials than that for vacancies ZVL around a dislocation. At a fixed temperature, MI is reasonably assumed to be constant. Therefore, the growth rate of the loops may be expressed as follows using the constant A to represent the coefficients in Eq. (6):
dR ¼ AC I dt
where C LI and C IF I indicate the concentrations of the interstitial atoms around the dislocation loop and the interface, respectively. C0 is the average defect concentration in the crystal grain. ULI and UIFI represent the interaction energy between the interstitial atoms and the loop and interface, respectively. k is Boltzmann’s constant and T is the thermodynamic temperature (K). In the general crystalline material the concentration of interstitial atoms around the dislocation loop can be described by Eq. (2). When a secondary phase precipitates in the matrix the concentration will decrease and can be described as:
0 U LI C LI ¼ C 0 C IF exp I kT
ð6Þ
ð2Þ
and
U IFI C IF I ¼ C 0 exp kT
dR ¼ aðZ IL bZ VL ÞMI C I dt
ð5Þ
This means that the interstitial concentration will decrease significantly if a large number of interfaces, which may be introduced by the nano-precipitates, form in the matrix of the materials. The rate of growth of the loop can be expressed as [44]:
ð7Þ
Here, the growth rate of the interstitial loop is controlled by the concentration C LI of the interstitial atoms around the dislocation loop. The interface between the precipitate and the matrix is an important factor that affects the concentration and the movement of point defects. This may provide vacancy-interstitial recombination centers that heal defects in the material. Therefore, the presence of high density precipitates will lead to a decrease in the growth rate of the dislocation loops by reducing the concentration of interstitial atoms around the dislocation loops and this may decrease the production of voids in these materials because of the annihilation of vacancy-interstitials at the interface. An analysis of the above-mentioned results reveal that the growth rate of dislocation loops was about 125.6 nm/dpa in the sample without the precipitate under 0.17 dpa irradiation. The growth rate was just 5.5 nm/dpa in the sample with precipitates. This means that the induced precipitates in this experiment greatly reduce the concentration of interstitial atoms around the dislocation loop. This leads to a significant decrease in the growth rate. In our experiment, no distinct void was found, even at a high temperature of 773 K with an irradiation dose of up to 2.13 dpa. Therefore, the nanoscale precipitates with high density that were induced by suitable pre-ion irradiation may lead to an improvement in the irradiation performance of these materials. This work thus provides a simple method to explore anti-irradiation in RAFM materials.
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5. Conclusions In addition to dislocation loops, irradiation with deuterium ions at 773 K induced the precipitation of a new phase rich in Cr in the Fe–10Cr model alloy. These nanoscale precipitates were arranged along the h1 0 0i direction with widths of 2–8 nm and lengths ranging from tens to hundreds of nm. The 1-D elongation of the precipitate is attributed to the distribution of strain fields around the precipitates. The results from a series of EDPs combined with diffraction calculations reveal that the phase crystallizes in the cubic space group and adopts a cube-onsystem. It belongs to the Pm3m cube orientation relationship with the Fe matrix and it has the same lattice parameters as Fe. Under subsequent 2 MV electron irradiation the precipitates are highly stable while the surrounding dislocation loops grow far more slowly compared with the loops in the specimen without ion irradiation induced precipitates. We argue that the distinct decrease in the growth rate of dislocation loops comes from the presence of a large number of fine precipitates. This is because the interfaces between the precipitate and the matrix may provide vacancy-interstitial recombination centers to heal defects in the material. We suggest that this precipitation at the nanoscale that is induced by deuterium ion irradiation may enhance the irradiation resistance of a material because of a reduction in the concentration of point defects. Acknowledgements This work is supported by the National Magnetic Confinement Fusion Program (Grant Nos. 2011GB108002, 2014GB104003 and 2014GB120001), the National Natural Science Foundation of China (Grant No. 51371031). The authors gratefully acknowledge the Research Center for Ultra-High Voltage Electron Microscopy at Osaka University, Japan. The help and suggestions from Professor H. Mori and Dr. K. Arakawa at Osaka University were highly appreciated. References [1] [2] [3] [4] [5]
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