Precipitates evolution during isothermal aging and its effect on tensile properties for an AFA alloy containing W and B elements

A 20Cr-25Ni-2.5Al alumina-forming austenitic alloy containing W and B elements was aged at 923 K for 5000 h, and the microstructure and tensile properties with different aging time were investigated. NiAl, σ and Laves were observed at grain boundaries (GBs) successively. The matrix was covered by NiAl and Laves with the extension of aging. The evolution of precipitates during aging contributed to the variation of tensile properties. Precipitation of nanosized NbC carbides within grains and σ phase at GBs led to a rapid increase in strength and a decrease in elongation for 500 h aging sample. In the later stage of aging, the coarsening of NiAl and Laves phases, as well as the reduction in dislocation density caused a decline in strength. The coarsening of precipitates upon aging time follows the Ostwald ripening theory. Due to its lower diffusion rate in austenite compared to Mo, W may accelerate the growth of Laves at GBs. Boron was mainly enriched in Laves instead of NiAl, NbC and σ phases after high temperature aging. The addition of W and B improved the precipitation strengthening of Laves, increasing the high temperature tensile strength after long term thermal aging. The difference in tensile properties between room temperature and 923 K is due to the ductile–brittle transition of NiAl. No σ phase was observed within grains even after 5000 h aging and elemental chromium particles occurred around Laves due to boron hindering the growth of σ.


Introduction
Alumina-forming austenitic (AFA) steel, developed by Oak Ridge National Laboratory (ORNL), is a kind of advanced stainless steel with excellent oxidation properties and high temperature creep resistance [1][2][3][4][5]. Protective Al 2 O 3 oxide film is formed on the surface of AFA alloys at 873-1123 K [4], making it more stable in harsh environment such as water vapor compared with traditional stainless steels with Cr 2 O 3 layer (such as Super304H and TP347HF) [6]. AFA alloy can be widely used in ultrasupercritical steamplants, chemical plants and other industries [7], and is also expected to serve as an advanced reactor cladding material [8]. AFAs mainly depend on precipitation strengthening and dislocation strengthening [9], both of which are related to the precipitated phases in the matrix. The precipitation in AFA is sequential [10,11]. The dissolution and coarsening of the second phases are essentially inevitable during high temperature exposure [12], thus long-term thermal aging is the main reason for mechanical properties degradation of austenitic heat resistant steels [13].
Some studies have confirmed the beneficial effects of W and B elements on the mechanical properties of stainless steels. Tarigan [14] proposed that the addition of boron could inhibit the local deformation near GBs and improve the creep resistance by forming smaller and more abundant Laves particles at high temperature, which is called ''grain boundary precipitation strengthening (GBPS)''. For (9-12)% Cr martensitic steels, boron is enriched at GBs to form M 23 (CB) 6 carbides, which can effectively reduce the coarsening of M 23 C 6 [15,16]. However, M 23 C 6 is not precipitated for most AFA alloys due to the stabilization of Nb or Ti [4]. Korean researchers used W to replace Mo to form a slowly coarsened Laves phase in AFA. Compared with Fe 2 Mo, the volume fraction of Fe 2 W increased and the coarsing rate became lower during aging [17][18][19][20], resulting in the improvement of the mechanical properties [21]. It is worth noting that studies on the aging behavior of AFAs containing both W and B are limited, and the distribution behavior of boron in precipitates is not clear.
At present, only a few studies on isothermal aging of AFAs used long duration experiments [22]. While in short-term high temperature exposure, there is almost no obvious coarsening of Laves and MC phase [23]. The coarsening rate of Laves may also not be constant in different aging periods with the consumption of Nb [9,10]. Therefore, for an AFA alloy containing W and trace B, the precipitation, growth and coarsening of second phases for 5000 h aging were carried out, and the influence of precipitates on tensile properties was also explored, which can deepen the understanding of the degradation of AFAs.

Material preparation
An alumina-forming austenitic steel containing W and B (AFA WB ) was adopted, and the chemical compositions are shown in Table 1. The ingot was prepared by vacuum induction melting. The casting ingot was heated up to 1473 K in the furnace, held for 2 h and then forged to a plate with a size of 100 9 100 9 30 mm (direction of forging), with the forging ratio of 3: 1. The forged ingot was homogenized for 1 h at 1473 K and then machined into small plates with a size of 50 9 20 9 15 mm by wire-electrode cutting. Isothermal aging for 5000 h at 923 K was carried out by using Muffle furnace. The microstructure characterization and tensile properties were tested for the samples aged for 0, 300, 500, 1000, 3000 and 5000 h, respectively.

Characterization
Microstructure characterization and phase identification were examined by using confocal laser scanning microscope (CLSM, OLYMPUS), scanning electron microscope (SEM, thermo scientific Aprep 2C) equipped with energy dispersive spectrometer (EDS, OXFORD ULTIM Max65) and CBS probe, transmission electron microscope (TEM, FEI Talos F200x) with EDS (Super X). Metallographic samples for SEM and optical microstructure (OM) analysis were prepared by wire-electrode cutting, ground with SiC papers up to 2000 grit, and were polished using non-crystallizing colloidal silica. The OM samples also need to be etched with 10 wt% oxalic acid. Image-Pro Plus software was used to calculate the average grain size and distributed information of precipitates including the average size and volume fraction. The thin foils for TEM tests were compared by using the twin jet method with a mixed solution of HClO 4 : CH 4 O = 10%: 90%. The thin foil containing GBs for TEM analysis was prepared by using focused ion beam (FIB).
In order to determine the effect of isothermal aging on the strength and plasticity, tensile tests were carried out at room temperature (RT, aged for 0, 1000, 3000, 5000 h) and at 923 K (aged for 0, 300, 500, 1000, 3000, and 5000 h). Rod-shaped tensile specimens ( Fig. 1) were adopted to carry out tensile tests by using the electronic universal testing machine (GNT100) and materials testing system (ELYF01) with an extension rate of 0.2 mm min -1 (1.2 mmÁmin -1 after yielding). The fracture morphology, precipitated phases and cracks in matrix were observed by using SEM.

Microstructure evolution of AFAWB
With the aging process, the second phases in AFA WB exhibited two stages successively: precipitation and coarsening. The precipitates in matrix after different aging time were shown in Fig. 2. After 1473 K solution treatment, the matrix was in form of austenite and only a small number of coarse particles with bright contrast that were identified as primary NbCs could be observed (Fig. 2a) [22]. The second phases, which played a key role in the mechanical properties of AFA WB , were precipitated in the subsequent aging process. After 300 h aging, precipitates with dark contrast were observed at GBs (Fig. 2b). Related studies suggest these particles may be NiAl or M 23 C 6 [17,18,22,24], and further characterization by using EDS point analisis confirmed they were the NiAl particles (Fig. 2k). After 500 h aging, Cr-rich r phase with gray contrast began to precipitate around the NiAl at GBs. The non-magnetic intermetallic compound r phase is rich in Fe and Cr with a bodycentered cubic structure [18], which has a higher coarsening rate in AFA and is harmful to ductility and toughness [9]. Due to the capture of C element by Nb or Ti, r phase usually precipitates at the early stage of aging [25]. The precipitation of NiAl and r led to the enrichment of W, Nb and other elements at GBs, thus helping to precipitate Laves (green arrows in Fig. 2c). There are usually three types of Laves: hexagonal MgZn 2 (C14), cubic MgCu 2 (C15) or hexagonal MgNi 2 (C36) [11], and generally a C14 type Fe 2 M-Laves (M=Mo, W, Nb) brittle phase precipitates in AFAs [26]. The precipitates at GBs caused the imbalance of elements distribution and gradually led to the precipitation from GBs into grains (Fig. 2e).
As for 1000 h aged sample, precipitates were mainly located near GBs, with no precipitation observed in most areas of the inner grains (Fig. 2d). These needle-like particles with dark contrast and the ones with bright contrast in grains were determined as NiAl and Laves by EDS, respectively. Some smaller Laves particles were spherical, while some that grew around NiAl were slightly larger with a flake or long strip shape. The morphology and identification of precipitates were characterized more accurately by using TEM. As shown in Fig. 3a, significant hindrance to dislocation movement by precipitates was observed. EDS result showed the uneven distribution of alloying elements in the matrix (Fig. 3c), in which Ni/Al rich area had an incoherent relationship with the austenite matrix (Fig. 3f). Fast Fourier transform (FFT) pattern conformed the phase is NiAl with BCC structure (Fig. 3g). Laves phase existed near NiAl (Fig. 3g). An interesting finding was that some particles rich in chromium were adjacent to NiAl and Laves, which were different from the same Cr-rich r phase at GBs ( Fig. 2e and b) by selected area electron diffraction (SAED) analysis (Fig. 3e). They were identified as elemental chromium particles with a BCC structure. The thin foil obtained by FIB was used to further analyze precipitation behavior at GBs (Fig. 4a). SAED patterns showed that precipitates located at GBs mainly consisted of r, NiAl and Laves ( Fig. 4b and c). The intermetallic compound NiAl has an B2-type ordered structure, and its diffraction resulted in superlattice pattern (marked by white circles in Fig. 4b). Secondary NbC particles existed within the Laves phase at GBs (yellow encircled area in Fig. 4c). The enrichment of niobium in Laves and its vicinity may provide an elemental source for the precipitation of NbCs.
When aging over 1000 h, the morphology and type of precipitates remained roughly unchanged, whether it was Laves, NiAl or r, but the quantity and distribution kept changing. When aging for 3000 h at 923 K, All parts of grains were covered with a large number of Laves and NiAl particles (Fig. 2g), and r phase continued to coarse at GBs (Fig. 2h), reaching micron size. The accompanying growth of NiAl and Laves was observed, and chromium particles free of Mo and W were also found to grow near NiAl and Laves (Fig. 5a). Nanosized NbCs of about 10 nm were also observed at the edge of Laves. The corresponding relationship between the redistribution of alloying elements and the precipitates was shown: Nb, Mo, W, Si (not shown in the figure) and B were enriched to form the Laves phase, with Ni and Al to form NiAl (Fig. 5b, c). Large r phase with gray contrast and black NiAl particles were found to grow together at GBs (Fig. 5d). A small amount of the Laves particles precipitated around r phase, but the size was significantly smaller than those near NiAl, which indicated that the growth of Laves around r may be inhibited.
When aging for 5000 h, the size of Laves and NiAl particles increased but the volume fraction basically did not change compared with 3000 h aging. The evolution of second phases transformed from precipitation to coarsening. The r particles at GBs were further coarsened, but it is noteworthy that no r was observed in grains even after 5000 h aging (Fig. 2j).
Grain boundaries were also covered by NiAl and Laves (Fig. 6a). EDS mapping showed in the intracrystalline region around the coarse r, Mo and W overlapped with Nb and B (yellow dotted box in Fig. 6b), indicating that boron is enriched in the Laves phase. Mo and W did not overlapped with Cr (white dotted box in Fig. 6b), so it shows a competitive relationship between Laves and r phase, both of which are formed by these elements' consumption. A string of NbC particles precipitated around dislocations in grains (blue dotted line in Fig. 6c). The FFT patterns of the NbC (encircled with red box A) and the matrix (encircled with blue box B) show that they have the following cube-on-cube orientation relationship: 110 This semi-coherent relationship between NbC and matrix has a lower interface energy [27], which causes NbC to coarsen slowly at high temperature. The more sluggish coarsening of NbC than other precipitates would be detailed in the Discussion section.
Dislocations in the matrix were also determined at different aging time by using TEM. After 300 h aging, a large number of dislocations, but a few second phases (due to the short aging time) were observed in grains (Fig. 7a). Nanosized NbC particles were observed following 500 h of aging. Upon further aging for 1000 h, precipitates with dark contrast were found in the matrix, with an average size of around 11 nm (blue arrows in Fig. 7b). These nanosized particles are rich in Nb but lack Mo, W, Si, and identified to be NbCs by FFT analysis (Fig. 7c). NbCs precipitated around dislocations and played a crucial role in hindering their movement, so as to improve high temperature strength of alloys. When aging for 5000 h, large Laves and NiAl particles as well as nanosized NbCs were observed within the grains.
NbC had an extremely gentle coarsening rate with an average size of 14 nm and maintained its thermodynamic stability. Dislocations clustered and entangled around Laves and NiAl, some forming dislocation loops, which made many NbCs lose the interaction with dislocations (blue arrows in Fig. 7d).
During the long term aging process at 923 K, dislocations that were introduced by forging gradually evolved and their density kept changing with migration and interaction with precipitates. According to the equation [28]: where n is the intersection count, l is the total length of the grid, and t is the thickness of the foil. Thus the dislocation density in AFA WB reached 6.04 9 10 12 m -2 , 8.63 9 10 13 m -2 , and 2.48 9 10 13 m -2 after aging for 300, 1000, and 5000 h, respectively. Metallic materials usually yield in the form of dislocation slip. An increase in dislocation density, although increasing the absolute number of dislocations, reduces movable dislocations and increases the stress required for dislocations to start slipping due to effective nailing of nanosized NbCs. Therefore, the variation of precipitated phases, together with dislocations evolution, is one of the important factors for the evolution of tensile properties during high temperature aging.

Tensile properties
The variation of tensile properties for different aging time at 923 K is shown in Fig. 8. The tensile curves only had a little change for 300 h aging sample and the as-solutionized one, including a slight decline in ultimate tensile strength (UTS), yield strength (YS) and strain after fracture. When aging for 500 h, the UTS of AFAWB increased rapidly from 416 Mpa (before aging) to the highest 684 MPa among all aged alloys. Then, the UTS maintained a high level of 655 MPa for 1000 h aging, and the YS increased to the highest 595 MPa rapidly from 208.5 MPa before aging (Fig. 8b). At the same time, the total strain decreased significantly, with the elongation value decreasing sharply from 44.75 to 12.75%, exhibiting an obvious embrittlement. As aging continued, the strength of the alloy decreased and the plasticity began to recover. After 3000 and 5000 h aging, the UTS of 485 and 449 MPa, respectively, showed a certain degree of degradation.
The fracture morphology and cracks after tensile tests at 923 K were investigated. Obvious dimple morphology was observed in the fracture of the assolutionized AFA (Fig. 9a), which was consistent with the high elongation value. This fracture characterized by dimples is a typical ductile fracture with high energy absorption process [29]. After 1000 h, both ductile fracture and cleavage fracture occurred ( Fig. 9(b1)), with dimples shallower in the depth direction. Correspondingly, the macro elongation decreased and the yielding-to-tensile ratio increased. Cleavage fractures usually occur along grain boundaries ( Fig. 9(b2)). After 3000 and 5000 h aging, the fracture exhibited dimple morphology again for 923 K tensile tests (Fig. 9c, d). Cracks in the matrix after fracture tended to be parallel to the loading direction, and the r phase and primary NbCs often acted as the priority for crack initiation (Fig. 9e, and f), which was consistent with earlier AFA research [30].
However, the variation in RT tensile properties (Fig. 10) differed from that at 923 K. The UTS continued to increase to 996 MPa (3000 h) and then slightly decreased to 975 MPa (5000 h), while the elongation remained at a very low value. The difference may be related to the B2-NiAl phase in the matrix. NiAl phase has a body-centered cubic structure with a ductile-to-brittle transition temperature (DBTT) within the range of 773-1073 K [4,24], so NiAl is the main strengthening phase at room temperature [31], whereas its strengthening effect may be lost at high temperature [32].

Kinetics for precipitation
The second phases would grow coarser continuously under high temperature aging. This coarsening behavior mainly depends on the growth of large particles by absorbing the mass of dissolved small and medium-sized precipitates [33]. The process is called Ostwald ripening and the coarsening rate of precipitates over time can be predicted by the following equation [21,23]: where t is aging time, d t and d 0 are the average particles size of the unaged sample and the aged sample for t, c is the interface energy between the precipitates and the matrix, V m is the molar volume fraction, D is the diffusion coefficient of solute, C m is the solute atomic fraction in equilibrium with the precipitates, R is the gas constant, T is the absolute temperature, and k is coarsening rate constant depending on c, V m , D, C m , R, and T. Considering that only a few particles were precipitated for 0-1000 h aging duration, statistics of precipitates for samples aged 1000-5000 h were carried out, and the curves reflecting cubic of average particle diameter as a function of aging time was drawn. As shown in Fig. 11, the d3-t curves roughly present a linear relationship for NiAl, Laves and NbC phases. According to Eq. (2), the slope is the coarsening rate constant k. The secondary NbC is abnormally stable with the coarsing rate (9.97 9 10 -32 ) nearly 3 orders of magnitude lower than that of NiAl (7 9 10 -29 ) and Laves (7.86 9 10 -29 ). This value of NbC is roughly equivalent to the one reported in the literature (923 K in P92 steel, 6.55 9 10 -32 [34]). Some studies have also confirmed that the extremely low coarsing rate of NbC at high temperature [35] is related to the slow diffusion coefficient of Nb in austenite and the low interface energy of NbC and matrix [22]. Nanoscale NbCs precipitated near dislocations (Fig. 7b), which generated an obvious pinning effect [36] and gave the alloy good strength.
Due to the earlier precipitation along GBs, the average size of Laves particles at GBs is slightly larger than that of intra-grain ones when aging for 1000 h. An interesting finding is that the coarseing rate of Laves at GBs is slightly lower than that within grains (Fig. 11), which is hypothesized to be related to W element. Laves is rich in Mo and W, and its precipitation and coarsening at GBs depend on the ''long-range'' diffusion of these elements in grains. Compared with Mo, W has a lower diffusion rate in the austenite matrix [17], thus slightly controlling the growth of Laves at GBs.

Distribution of B and its effect on precipitates
Boron is segregated and enriched at grain boundaries after annealing. B can enhance the precipitation of Laves phase at GBs, but inhibit its coarsening [37]. Tarigan [14] found that a Fe-20Cr-30Ni-2Nb alloy containing 0.03% B could significantly increase the creep life, with more compact and fine Laves particles in crept specimen. However, no relevant studies on the specific distribution of boron in precipitates and the resulting effects on the mechanical properties for AFA steels have reported.
According to the microstructure analysis for aged AFA WB , the distribution behavior of B in the precipitated phases can be summarized. Enrichment of B occurs in the Laves phase both at GBs (Fig. 12b) and within grains (Fig. 6d, Fig. 12c), but does not in NiAl (Fig. 6b, Fig. 12b), r phase (Fig. 6b) and NbCs (Fig. 6d). A recent study has also confirmed the segragation of B at the interface between Fe2W and matrix for a 3Co-3W-9Cr martensitic steel [38]. This preferred distribution existed for aging both 1000 h ( Fig. 12) and 5000 h (Fig. 6), indicating the stability of boron's distribution in Laves.
For the AFA WB in present work, the addition of boron may cause some changes of high temperature precipitation. In traditional AFA alloys (without W and B), Laves phase tends to precipitate near NiAl or r phase [22]. This associated precipitation behavior has also been observed in AFAs containing W [18]. However, The elementary chromium particles instead of r phase were observed around NiAl and Laves precipitated earlier in AFA WB , whether aging for 1000 h (Fig. 3b), 3000 h (Fig. 5a) or 5000 h (Fig. 12a). This was clearly shown in the EDS mapping for 5000 h aging (Fig. 12d). Laves (rich in W and Mo) tended to precipitate and grow adjacent to NiAl (rich in Al), while chromium particles free of W and Mo were precipitated near Laves instead of r. Only some Laves particles, yet with significantly smaller size, occur near r phase at GBs (Fig. 2j, Fig. 5c). This suggests that r could hardly grow through the dissolution of Laves, whereas Laves can precipitate by consuming r phase. According to the microscopic characterization of aged samples, there are mainly three types of associated precipitation behavior in the late stage of aging (Fig. 13a): NiAl ? Laves ? Cr within the grain (I), NiAl ? Laves ? Cr at GB (II) and NiAl ? early precipited r ? small size Laves at GB (III). Boron may somehow inhibit the diffusion of W and Mo in the Laves phase to the matrix and make it difficult for them to combine with the surrounding Cr (Fig. 13b), thus inhibiting the growth of r phase.

Effect of precipitates on tensile properties
After high temperature aging at 923 K, the mechanical properties of AFA WB exhibited initial strengthening followed by subsequent degradation (Fig. 8).
The reason for the change of tensile strength can be expressed by Eq. (3) [39]: where r represents the yield strength of materials and to some extent also can reflect the tensile strength. r 0 is the inherent strength of single crystal pure matrix, r ss is the solution strengthening from alloying elements, r gb is the grain boundary strengthening, r d is the dislocation strengthening and r p is the precipitation strengthening. r 0 should be the same for different aging time. r gb is closely related to the grain size. The relationship between the grain size and strength can be expressed by the Hall-Petch formula [40]: where d is the average grain diameter, r f0 is the friction stress acting on dislocations, k is a yield constant. As it is shown in Table 2, The average grain size remained stable before 1000 h aging, after which it began to increase. Grain boundary migration occurs at high temperature, which leads to the grain size change and can be hindered by precipitated phases at GBs. The nailing force of the precipitates is expressed by [10]: where K is a constant related to the material, f and r are the volume fraction and average radius of the precipitates, respectively. This means that the higher the number of second phases and the smaller the size, the better the nailing effect on the grain boundaries. Compared with the as-solutionized sample, UTS and YS after aging for 300 h did not increase, but slightly decreased. Yet a certain amount of nanosized NbCs had not precipitated within the grains (Fig. 7a), so their effective resistance to dislocation motion could not be formed. The NiAl phase precipitated at GBs (Fig. 2b), but there was no significant strengthening effect due to its tough-brittle transition. After 500 h aging, precipitation of NbC was observed in the matrix (Fig. 7b), which coincided with a significant increase in high temperature strength. The smaller the size and the greater the quantity of precipitated phases are, the greater the pinning to dislocations and the higher the strength will be, which is given by the Orowan mechanism [41]: where r or is the critical stress to bypass dislocations, G is the shear modulus, b is Burgers vector, k is the planar spacing between particles. So, the nanoscale particles can effectively hinder dislocations. The precipitation of secondary nanoscale NbC has remarkable effects on the high temperature strength and creep resistance of AFAs [4,21,32]. On the other hand, precipitation of r phase at GBs was observed (Fig. 2c), which could significantly enhance the strength of AFAs [42,43]. However, this also leads to brittle fracture and provides a path for crack propagation [44], resulting in decreased toughness and ductility. Therefore, the precipitation of r at GBs caused an increase in r p , and the NbCs in grains led to an increase in r d . These combined effects ultimately caused significant changes in tensile properties of the 500 h aged sample.
As the evolution of precipitates entered the Ostwald ripening stage, the strength at 923 K slowly decreased, while the elongation began to rise. After 5000 h aging, the average size of NiAl and Laves reached 130 and 109 nm, respectively. Compared to NbC of only 11.2 nm, r or obtained by NiAl and Laves began to decrease. The toughness of NiAl above DBTT also makes the dislocations less obstructed and easier to move. Dislocation strengthening r d can be calculated by [45]: where M = 3.06 is the mean orientation factor for the fcc matrix, a is a constant of * 0.2 for fcc alloys [45] G and b for AFA are 80 GPa and 2.5 9 10 -4 lm, respectively [46], and q is the dislocation density. According to Eq. (7), The r d of AFAWB reaches the highest of 113.7 MPa after 1000 h aging and then decreases gradually. At the same time, the content of alloying elements such as C, W, Mo, Nb, Al and Si in the matrix gradually decreases with second phases precipitating and growing, and so does r ss . According to Eq. (5), the continuous coarseing of Laves, NiAl and r phases at GBs (Fig. 2e, h, j) leads to the weakening of the hindrance to grain growth (Table 2). Therefore, the simultaneous decrease of r d , r ss , r p , and r gb results in a reduction of macroscopic strength and an increase in plasticity during aging for more than 1000 h. The addition of W and B into AFA alloys induces distinct degradation behavior at high temperatures.
A comparison of the tensile strength between AFA WB and other AFA alloys mentioned in literature without W and B is presented in Fig. 14 The UTS and YS of AFA alloys at room temperature increase and the elongation decreases accordingly with aging. Although the tensile strength of as-solutionized AFA WB is not outstanding, it exhibits more significant precipitation strengthening (more substantial increase in UTS, Fig. 14a) after thermal aging. Interestingly, the high temperature tensile properties of these alloys remain largely unaffected by the aging process-either exhibiting minimal changes [4] or displaying degradation in mechanical performance characterized by a gradual decline in strength over time (Fig. 14b). This distinct tensile behavior observed at elevated temperatures can be mainly attributed to DBTT of the NiAl phase [12,31].
Unlike AFA alloys without W and B, the high temperature strength of AFA WB is significantly increased after aging. As previously mentioned, the increase in strength of 500 h aging sample is mainly attributed to the precipitation of nanoscale NbCs in grains and r phase at GBs. The subsequent decrease in strength is closely linked to the coarsening of second phases. It is noteworthy that the high temperature strength of 5000 h aging sample remains higher than that of the as-solutionized sample. Considering the weakened strengthening effect of NiAl at Figure 11 The cubic of average particle diameter of AFA WB as a function of aging time during isothermal aging (slope is coarsening rate constant k).
high temperature, the enhanced strength mainly originates from the improved Laves strengthening. Compared to Mo, W has a higher affinity for Laves, which makes the Laves phase more stable [10]. The slower diffusion rate of W effectively suppresses the coarsening of Laves at GBs (Fig. 11). On the other hand, the enrichment of B in the Laves phase effectively impedes its transition into r phase (Fig. 13). Therefore, W and B elements make AFAs perform better in terms of high temperature tensile properties during prolonged aging.
It should be noted that the r phase along GBs has great harm to mechanical properties and is the main location for cracking initiation under external force loading (Fig. 9f), resulting in rapid fracture failure of materials with low plastic deformation, which is the microscopic reason of material embrittlement. The adoption of B makes it difficult to form r phase around the Laves (Fig. 12d), thus reducing the harm of r phase to a certain extent. However, Al promotes the growth of r by precipitating the NiAl phase [24]. Therefore, to further inhibit the ripening of r in AFAs, it is a key to obtain a better high temperature aging performance.

Conclusions
An AFA alloy containing W and B (AFA WB ) was aged for 5000 h at 923 K, and precipitation behavior and its effect on tensile properties during aging were studied. The main conclusions are as follows: (1) During aging at 923 K, NiAl, r and Laves were observed at GBs successively, and then NiAl and Laves were precipitated at the intergranular region near GBs, to cover the entire grain with aging continuing. No in-grain r phase was observed even after 5000 h aging. Nanosized NbCs mainly precipitated around dislocations and Laves particles. (2) The coarsening of precipitates upon isothermal aging follows the Ostwald ripening equation.
The coarsing rate constants (k) of secondary NbCs (9.97 9 10 -32 ) is nearly 3 orders of magnitude lower than that of Lave phase and NiAl phase (7 9 10 -29 and 7.86 9 10 -29 , respectively) at 923 K. The lower diffusion rate of W in Figure 13 a Schematic diagram of distribution of precipitated phases during 923 K isothermal aging; b illustration of the suppression of r growth by element B.  Figure 14 Comparison of UTS between AFA WB and AFAs without W and B at a room temperature; b elevated temperature.
austenite than that of Mo may be the reason for slower coarsing rate of Laves at GBs than in grains. (3) The distribution behavior of boron in precipitates for AFA is described for the first time. B is mainly enriched in Laves instead of NiAl, NbC and r phase after thermal aging. The Cr-rich phase formed as elemental chromium particles around the Laves phase, which differs from traditional W and B-free AFAs precipitation. The addition of B may inhibit the diffusion of W and Mo in Laves phase towards adjacent Crrich regions, thereby hindering the growth of r phase. (4) The resistance of NbC carbides to dislocations and the precipitation of r phase at GBs led to the rapid increase of 923 K strength and the corresponding decrease of elongation of 500 h aging sample. At the later stage of aging, a slow decline in 923 K strength and a rise in plastic were observed, which was related to the ripening of NiAl and Laves and the decrease of dislocation density. The ductile-brittle transition behavior of NiAl was responsible for the difference in tensile properties between room temperature and 923 K. The addition of W and B enhanced the precipitation strengthening of Laves, thus increasing the high temperature tensile strength of AFA after thermal aging.