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The effect of boron on B2 compounds
The effect of boron on B2 compounds
7.1
7
NiAl
As indicated earlier ordered B2 NiAl has excellent properties among them low density, good oxidation and corrosion resistance and high melting point, therefore a possible high temperature structural material for industrial applications. However, the inherent brittleness at low temperature limited its use for aeroengine manufacturing. The grain boundaries were considered the weak region of NiAl where intergranular cohesion is not sufficiently strong and the failure by fracture on the application of load is considered to be intergranular. The addition of proper alloying elements is an effective method to improve the mechanical properties of intermetallic compounds such as NiAl. Such additives not only strengthen the matrix, but also significantly enhance the intergranular cohesion due to their segregation in grain boundaries. B is considered an excellent strengthener in various intermetallic compounds among them NiAl also. Experiments were performed to evaluate the effectiveness of B to suppress grain boundary fracture in NiAl and correlate it with grain boundary segregation and the improvement of the ductility. Table 7.1 shows the room temperature properties with and without B additions. C and Be are included in the table, since first principle calculations indicated that C is better than B in improving cohesive strength and Be was found to have a beneficial effect in the case of Ni3Al. One can observe in the table that the UTS is increased, both in the partially relaxed and relaxed stage. The expectation of improved ductility was not fulfilled when 300 ppm B was added, but grain boundary fracture was suppressed almost completely. Actually it is claimed that the B doped NiAl has a lower ductility than the undoped NiAl. In the experiments performed not showing improved ductility, can be explained by the fact that the high tensile strength in general is on the expense of elongation, namely the variation of strength is found to be opposite to the ductility. Increase in strength generally is not accompanied with increase in ductility. The addition of 300 ppm B, however, suppresses grain boundary fracture. Fig. 7.1 shows room-temperature fracture surface of B free and B containing NiAl. In order to determine if there was any change in the composition at grain boundaries relative to the bulk due to B additions plots of the peak height ratios (PHRs), namely that of Ni (848 eV)/Al (1396 eV) for both intergranular and transgranular areas was made for the alloys. This is seen in Fig. 7.2. Clearly the elemental peak heights were determined by Auger spectra. The Ni and Al peak spectra in the B doped NiAl at grain boundaries are higher than are without doping, an indication for B segregation in grain boundaries. No impurities Basic Compounds for Superalloys. DOI: https://doi.org/10.1016/B978-0-12-816133-3.00007-8 © 2018 Elsevier Inc. All rights reserved.
254
Basic Compounds for Superalloys
Room-temperature tensile properties of NiAl microalloyed with B, C, and Be (George and Liu, 1990)
Table 7.1
Alloy composition (at.%)
Microstructure (heat treatment)
Yield strength (MPa)
Ultimate tensile strength (MPa)
Tensile elongationa (%)
Ni50Al Ni50Al NiAl 1 300 ppm Bb NiAl 1 300 ppm Bb NiAl 1 300 ppm Cb NiAl 1 300 ppm Cb NiAl 1 500 ppm Beb NiAl 1 500 ppm Beb
Partial rex. (1 h/500 C) Rex. (1 h/805 C) Partial rex. (1 h/800 C) Rex. (1 h/1000 C) Partial rex. (1 h/500 C) Rex. (1 h/1100 C) Partial rex. (1 h/950 C) Rex. (1 h/1050 C)
198 154
262 229 400 329 429 336 268 307
2.0 2.2
c c c c
188 178
c c c c
1.9 3.0
a
Elongation measured from a strip chart. Weight parts per million. Fracture prior to macroscopic yielding. Source: With kind permission of Cambridge University Press.
b c
Figure 7.1 Room-temperature fracture surfaces investigated in this study (George and Liu, 1990). Source: With kind permission of Cambridge University Press.
were detected in the stoichiometric NiAl which could possibly explain the intergranular fracture, an indication that NiAl is intrinsically brittle. Of interest is to note that the Ni-rich B2 NiAl shows not only martensitic transformation, but also perfect pseudoelasticity. The 37.5 at.% Al alloy is off-stoichiometry, but still having a B2 structure. B is added because it is one of the most effective microalloying elements improving the ductility and suppressing the intergranular fracture. It changes the Ms temperature but to a smaller degree the Mf temperature and influences the reverse transformation temperature. Also the B addition seems to restrict grain growth to some extent and the 37.5 at.% AlB has a smaller grain size. X-ray diffraction indicates that both alloys are ordered B2 structures. The phase transitions was evaluated from expansion vs temperature curves, shown for the 37.5 at.% Al and the 37.5 at.% AlB NiAl alloys in Fig. 7.3.
The effect of boron on B2 compounds
255
NiAl
NiAl + 300 ppm B 4.0 IG TG
3.5
Peak height ratio (Ni/Al)
Peak height ratio (Ni/Al)
4.0
3.0 2.5 2.0 1.5 1.0
3.0 2.5 2.0 1.5 1.0
0
5
10
15
20
25
IG TG
3.5
0
10
Analysis No.
20
30
Analysis No.
Figure 7.2 Ni/Al PHRs for various intergranular and transgranular areas in the alloys analyzed by Auger electron spectroscopy (George and Liu, 1990). Source: With kind permission of Cambridge University Press.
(A)
As Linear expansion
Mf Af Ms
(B) Mf.As
Af Ms –200
–100 0 Temperature (°C)
100
Figure 7.3 (A) Linear expansion-temperature curves of the Ni-37.5 at.% Al specimen and (B) the Ni-37.5 at.% AlB specimen (Xie et al., 1993). Source: With kind permission of Elsevier.
Note that B additions modify the Ms and the Af temperatures, but almost no change has occurred in the Mf and As temperatures (Table 7.2). The As refer to the reverse transition (in FeC phase diagram, it refers to Austenite). Fracture occurs readily by a very small force without B addition due to the extreme brittleness of the annealed Ni 37.5 at.% alloy. The intergranular fracture of the
256
Basic Compounds for Superalloys
Phase transformation temperatures of the alloys (Xie et al., 1993)
Table 7.2
Alloy
Ms ( C)
Mf( C)
As ( C)
Af( C)
Ni37.5 at.% Al Ni37.5 at.% AlB
2 118 2 84
2 129 2 129
2 121 2 129
2 111 2 75
Source: With kind permission of Elsevier.
Figure 7.4 (A) Morphologies of the fracture surface of the Ni-37.5 at.% Al specimen and (B) the Ni-37.5 at.% AlB specimen (Xie et al., 1993). Source: With kind permission of Elsevier. 700
Stress (MPa)
600 (A)
500
(B)
400 300 200 100
0.25 % 0
0 Strain
Figure 7.5 Compressive stressstrain curves of the Ni-37.5 at.% AlB specimens (Xie et al., 1993). Source: With kind permission of Elsevier.
unalloyed Ni 37.5 at.% Al is seen in Fig. 7.4A and some ductility was observed in the 37.5 at.% AlB NiAl, shown for comparison in Fig. 7.4B. The transformation temperatures of the alloys marked in Fig. 7.3 are shown in Table 7.2. Compressive stressstrain curves of the B doped alloy are illustrated in Fig. 7.5.
The effect of boron on B2 compounds
257
There is a definite yield point at 0.9% strain, as marked by the arrow and a total strain of 3.2% is obtained before fracture. The strain is recovered completely on unloading of the specimen compressed to 2.8% strain as seen in Fig. 7.5B. This observation of perfect pseudoelasticity is similar to the observations in shape memory alloys. Thus a stress induced martensitic transformation is taking place during loading and its reverse deformation occurs on unloading. B modifies the Ms temperature and induces a complete pseudoelasticity strain of 2.8% at room temperature.
7.2
FeAl
As in other B2 compounds additions of small amounts of B may change the fracture mode from intergranular to transgranular and induce ductility in these alloys at room temperature. Room temperature ductility is clearly very important, allowing machinability in inherently brittle materials, by alloying with small amounts of an additive. In the following information on FeAl alloys of different Al and B concentrations is presented. Due to the large homogeneity range showing an extended phase field in FeAl at the location of B2, it is possible to obtain iron aluminides of off-stoichiometric compositions, where the deficiency of atomic species leads to constitutional vacancies formation in order to maintain charge neutrality. It is known that thermal vacancies harden aluminides and quenched in vacancies can be relieved by annealing at some temperature. The concentration of vacancies was reported to increase with Al content. The hardening of FeAl depends on the substitutional element. In this book the concern is the effect of B as the additional element. The reports on the effect of B on the strengthening of FeAl are not uniform and these are considered in the following text. B in the amounts of 0.05, 0.1, 0.15, 0.3, 0.5, 1, 2, and 4 at.% were added to the base Fe45 at.% Al alloy. As cast Rockwell hardnesses of FeAl with B additions were measured prior to heat treatment of each specimen at 950 C for 2 h in air followed by a vacancy annealing of 400 C for 120 h in air. Hardnesses were measured also after vacancy annealing. The effect of B on the Rockwell hardnesses is shown in Fig. 7.6 indicating the as-cast and heat treated hardnesses. The hardness of the as cast specimens is the highest at all B concentrations. However, the second anneal at 950 C for 2 h and 400 C for 120 h increased the hardnesses at all B concentration but to a level still lower than the as cast specimens. Also note that the hardness of B containing alloys at all B levels is lower than the unalloyed FeAl. It is interesting to note that the 0.2% yield strength is significantly higher than the unalloyed FeAl as seen in Fig. 7.7. B doped alloys show only a very small decrease in yield strength after the vacancy-annealing treatment. The oil-dipped specimens indicated in Fig. 7.7 was intended to eliminate possible hydrogen embrittlement and notch effects to which FeAl tested in air is very sensitive. No meaningful difference in strength between the oil dipped and undipped specimens is observed. The effect of B on tensile elongation is seen in Fig. 7.8. The tensile elongation is low in both, the
258
Basic Compounds for Superalloys
80 As Cast 1st-950°C/2 h 1st-400°C/120 h 2nd-950°C/2 h 2nd-400°C/120 h
Hardness (Rc)
60
40
20
0
0
0.05
0.1 0.15 % Boron
0.3
0.5
Figure 7.6 Effect of boron content on the hardnesses of Fe45 at.% Al alloys in the as-cast condition and after annealing at 950 C for 2 h, then air cooling, and subsequent aging at 400 C for 120 h (Deevi et al., 1997).
0.2% Yield strength (MPa)
750 950°C/2 h Air cool 400°C/120 h Air cool 950°C/2 h Air cool-Oil Dip 400°C/120 h Air cool-Oil Dip
500
250
0 Fe–-Al
0.05 0.1 Boron content (at.%)
0.15
Figure 7.7 Effect of boron content on room-temperature yield strengths of Fe45 at.% Al alloys (Deevi et al., 1997).
B-doped and B-free specimens, but it increases after the vacancy annealing treatment (at 400 C for 120 h) to 0.46%. The alloys with 0.1 at.% B had the highest elongation of 1.9% after vacancy annealing treatment. The microstructures of annealed and vacancy annealed treated samples were quite similar, but a small grain size reduction was observed in the B-doped specimens in the range of 0.050.15 at.% B. The increase in room temperature properties such as yield strength and UTS by B addition is attributed to the grain boundary strengthening (Table 7.3).
The effect of boron on B2 compounds
259
Total elongation (%)
3 950°C/2 h Air cool 400°C/120 h Air cool 950°C/2 h Air cool-Oil Dip 400°C/120 h Air cool-Oil Dip 2
1
0
0
0.05 0.1 Boron content (at.%)
0.15
Figure 7.8 Effect of boron content on the room-temperature tensile elongations of Fe45 at.% Al alloys (Deevi et al., 1997).
B is claimed to segregate to grain boundaries which is considered by some to be the cause of the strengthening effect. The room temperature fracture mode changes at high temperatures from intergranular to transgranular, namely from brittle to ductile, as observed in both alloys the alloy of Fe46.1 at.% Al and the Fe46.1 at.% Al-0.78 at.% B. The engineering stressstrain curves for these two alloys at different temperatures are shown in Fig. 7.9. The tensile properties are listed in Table 7.2. The grains were equiaxed after the homogenization treatment in both alloys indicating recrystallization as seen in the microstructure of Fig. 7.10. The second-phase particles in the FeAlB specimens are suspected to be borides. In the FeAl specimens—which were totally brittle up to 600K—a significant ductility of 8% elongation in one specimen and 55% plastic elongation to fracture in another were measured. Disappointedly the FeAlB alloy remains brittle at 300K, but shows some ductility at 560K. Thus only a modest increase in ductility is the result of B addition. However a significant increase in the strength of FeAl is obtained by the addition of B. Also the addition of B changes the fracture mode from intergranular to transgranular, but the alloy remains brittle. The fracture surfaces of these two alloys are shown in Fig. 7.11 and as seen, at 300K, the FeAl specimen is totally intergranular, whereas the FeAlB specimen displays totally transgranular cleavage fracture. At 640K—just above the ductile-to-brittle transition temperature—the intergranular mode for FeAl and transgranular mode for FeAlB is maintained and still observed but some ductility is present. The FeAl specimen which showed a large ductility at 640K is characterized by dimple rupture features as seen in Fig. 7.12. The dislocation structures of homogenized FeAl and FeAlB specimens are illustrated in Fig. 7.13. The dislocations were introduced during the hot extrusion process and were sessile in their nature to be removed by heat treatment. They had
Results of tensile tests on FeAl and FeAlB as a function of temperature (constant strain rate, about 3 3 1023s21) (Crimp and Vedula, 1986)
Table 7.3
Temperature (K)
FeAl Yield strengtha (klbf in22)
300 560 600 640 (1) 640 (2) a
Ultimate tensile strengthb (klbf in22)
Fracture stressc (klbf in22)
FeAlB Elongationd (%)
Yield strengtha (klbf in22)
Ultimate tensile strengthb (klbf in22)
77 80 89 71
92 85
8 55
137 120 116
Fracture stressc (klbf in22) 141 143 130
129
Elongationd (%)
,1 2.5 11
0.2% offset yield strength for specimens which exhibited plastic yielding. Ultimate tensile strength for specimens which showed maximum load followed by necking elongation. Stress at which brittle fracture occurred before yielding for brittle specimens or stress at which ductile fracture occurred before necking elongation for ductile specimens. d Percentage elongation to failure in a gauge length of 3 cm for ductile specimens. Source: With kind permission of Elsevier. b c
The effect of boron on B2 compounds
261
–2
klbf in
MPa
160 1000
Stress
120 640 k 600 k 300 k
80
640 k 0.55 Strain at 500 fracture
40
0 0
0.04
0.08
(A)
0.12 Strain
0.16
0.20
klbf in–2
MPa
160 300 k 560 k 600 k
1000 640 k
Stress
120
80
500
40
0 (B)
0.04
0.08
0.12 Strain
0.16
0.20
Figure 7.9 Comparison of engineering stressstrain curves for (A) FeAl and (B) FeAlB alloy (Crimp and Vedula, 1986). Source: With kind permission of Elsevier.
definite orientations and the dislocations consisted of sharp jogs without evidence of faults or superlattice dislocations. Addition of B had a noticeable effect on the substructure of the dislocations by decreasing the tendency of having a specific orientation, making the jogs smoother and in inducing stacking faults formation. Because of B segregation in grain boundaries a strengthening of boundaries relative to the matrix was observed, as a consequence of modifying the fracture mode from intergranular to transgranular. However, no measurable ductility at room
262
Basic Compounds for Superalloys
Figure 7.10 Comparison of optical micrographs of the two alloys, illustrating recrystallized grains and prior particle boundaries in both alloys: (A) FeAl, longitudinal, (B) FeAl, transverse, (C) FeAlB, longitudinal, and (D) FeAlB, transverse (Crimp and Vedula, 1986). Source: With kind permission of Elsevier.
temperature was observed despite the change occurring in the fracture mode. Despite the relatively high level of B addition (0.78 at.%) which is beyond the amount required for strengthening grain boundary cohesion, it is speculated that cross-slip may contribute to early crack formation before plastic deformation could take place across grain boundaries. The concept is thus associated with the stress required to initiate plastic flow. There is a competition between this stress and the one required for brittle fracture either intergranular or transgranular. If the stress required to initiate brittle fracture (intergranular or transgranular) is less than that required to cause plastic flow, the specimen will break in a brittle manner, whereas in the opposite case (stress for plastic flow lower than the brittle fracture stress) as it is at higher temperatures the specimen exhibits ductility. The addition of B modifies the dislocation structure by forming stacking faults, which without B additions is a material without faulted structure. The presence of the stacking faults makes slip more difficult resulting in the increase in the matrix strength. The possibility of solid solution strengthening by B and particle strengthening are not ruled out. Thus
The effect of boron on B2 compounds
263
Figure 7.11 Comparison of scanning electron micrographs of the fracture surfaces for the two alloys at two different temperatures: (A) FeAl, 300K; (B) FeAl, 640K; (C) FeAlB, 300K; (D) FeAlB, 640K (Crimp and Vedula, 1986). Source: With kind permission of Elsevier.
a twofold strengthening by B additions should be noted: (a) B segregation at grain boundaries and modifying at low temperatures the fracture mode from intergranular to transgranular and (b) matrix strengthening by the formation of stacking faults making cross-slip more difficult. A different interpretation of the effect of B addition involves the influence of vacancies in the FeAl. Large grain alloys (B400 μm) containing 40, 46, or 50 at.% Al with additions of 80200 appm (atomic parts per million) are considered. A two-stage heat treatment was applied to these alloys: (a) to induce thermal vacancies at 950 C/1 h to obtain a maximum concentration of vacancies followed by air quenching and (b) 400 C/24 h for the elimination of excess thermal vacancies. Experimental results showing B concentration was evaluated either by the McLean or the Fowler’s model given as McLean
ΔGsegr cj cv 2 cj 5 exp 2 1 2 cv kT c0j
(7.1)
264
Basic Compounds for Superalloys
Figure 7.12 Scanning electron microscopy fracture surface of FeAl alloys showing (A) large ductility and (B) typical ductile fracture at 640K (Crimp and Vedula, 1986). Source: With kind permission of Elsevier.
Fowler
2 ΔGsegr 2 zw:cj =c0j cj cv 2 c 5 exp j 1 2 cv kT c0j
! (7.2)
where cj is the intergranular segregation of the segregating atoms,c0j is the saturation level at grain boundaries (considered to be a monolayer), cv is the bulk concentration representing the solubility limit, ΔGsegr is the energy of segregation and zw is an interaction factor. According to the sign of the zw parameter, the segregation may be limited (zw . 0-repulsive interactions) or enhanced (zw , 0-attractive interactions). The effect of bulk concentration, cv on intergranular segregation, cj is illustrated in Fig. 7.14. The intergranular concentration cj as a function of temperature is seen in Fig. 7.15. No significant influence by heat treatment in the temperature range 300950 C could be observed and the segregation level is almost the same as in the quenched alloy. This observation does not support an equilibrium segregation model. Both the equilibrium model of McLean and Fowler predict a strong decrease of the intergranular model with temperature [Eqs. (7.1)and (7.2)].
The effect of boron on B2 compounds
265
Figure 7.13 Transmission electron micrographs illustrating the major differences in the substructure of the two alloys (A) FeAl and (B) FeAlB (Crimp and Vedula, 1986). Source: With kind permission of Elsevier.
Because of the large difference in the energies of formation and migration, (0.7 and 1.7 eV/atom, respectively) the quenched alloys are able to retain the high concentration of thermal vacancies up to 0.2 at.% in Fe40Al and even up to 2 at.% in Fe50Al alloys (Paris). These vacancies—in excess at room temperature—may be eliminated at a low temperature heat treatment. The elimination of the vacancies can be evaluated from density measurements. This is shown for pure and B doped FeAl of 40, 46, and 50 at.% Al in Figs. 7.167.18, respectively. As indicated earlier the alloys were air quenched from 950 C. Under the assumption that one controlling mechanism of vacancy elimination is dominant, the decrease in vacancy concentration dN/dN0 may be expressed as dN K 5 2 exp dN0 kT
(7.3)
266
Basic Compounds for Superalloys
18 Exp (max) Fowler McLean
16 14 Cj (at%)
12 10 8 6 4 2 0 500
0
1000 Cv (appm)
1500
2000
Figure 7.14 Effect of the bulk boron concentration (cv) on its intergranular concentration (cj) in Fe40Al alloys. Heat treatment: 950 C/1 h air quench 1 400 C/24 h (Fraczkiewicz et al., 1998). Source: With kind permission of Elsevier. 9 8 7 Cj (at%)
6 5 4 3 2 1 0 200
300
400
500 600 700 Temperature (°C)
800
900
1000
Figure 7.15 Effect of temperature on the intergranular concentration of boron. Fe45Al 1 400 appm B alloy; annealing during 24 h (Fraczkiewicz et al., 1998). Source: With kind permission of Elsevier.
K is the reaction constant. The ratio of Kdoped/Kpure as a function of bulk B concentration is given in Fig. 7.19. Clearly the time of annealing influences the concentration of the vacancies as can be seen in Fig. 7.20 when it is carried out at 400 C. The acceleration of vacancy elimination kinetics in the presence of boron can be explained by the formation of boron-vacancy complexes, migrating to the grain boundaries faster than the desegregated B returning back to the matrix. Nevertheless, still an equilibrium segregation of B at a low level remained. A new dislocation structure is formed following the elimination of thermal vacancies. In the pure FeAl, vacancy elimination following quenching, is associated already at the start with dislocation loops creation with h111i Burgers vector. But
The effect of boron on B2 compounds
267
Atomic fraction of vacancies (%)
0,25 Fe–40Al–200B Fe–40Al–40B
0,2
Fe–40Al
0,15 0,1 0,05 0 0
2
4
6 8 10 Annealing time (h)
12
14
16
Figure 7.16 Kinetics of quenched-in vacancies elimination during an isothermal annealing at 380 C. Dilatometric study of Fe40Al alloys, pure or B-doped (Fraczkiewicz et al., 1998). Source: With kind permission of Elsevier.
Atomic fraction of vacancies (%)
0,45 Fe–46Al–200B Fe–46Al
0,4 0,35 0,3 0,25 0,2 0,15 0,1 0,05 0
0
5
10
15 20 Annealing time (h)
25
30
35
Figure 7.17 Kinetics of quenched-in vacancies elimination during isothermal annealing at 380 C. Dilatometric study of Fe46Al alloys, pure or B-doped (Fraczkiewicz et al., 1998). Source: With kind permission of Elsevier.
after annealing the homogeneous h100i dislocation network is obtained. This can be seen in Fig. 7.21. In B doped FeAl the dislocation structure is different as shown in Fig. 7.22. Vacancy elimination starts faster than in pure FeAl. As a result the dislocation network due to vacancies elimination is almost the same in the quenched state as in the annealed samples for 24 h at 400 C (Fig. 7.22A); it is regular but heterogeneous. The cell walls are B3 μm in size with only a few individual dislocations (namely high dislocation density walls). After 400 C annealing for a week a homogeneous structure similar to the pure alloy is obtained as seen in Fig. 7.22B. Despite the acceleration macroscopic effect of vacancy elimination by B addition
268
Basic Compounds for Superalloys
Atomic fraction of vacancies (%)
0,8 Fe–50Al–200B Fe–50Al Fe–50Al–800B
0,7 0,6 0,5 0,4 0,3 0,2 0,1 0 0
10
20 30 40 Annealing time (h)
50
60
Kdoped/Kpure
Figure 7.18 Kinetics of quenched-in vacancies elimination during isothermal annealing at 380 C. Dilatometric study of Fe50Al alloys, pure or B-doped (Fraczkiewicz et al., 1998). Source: With kind permission of Elsevier. 4,5 4 3,5 3 2,5 2 1,5 1 0,5 0
Fe–40AI Fe–46AI Fe–50AI
0
0,02
0,04
0,06 0,08 Cv/No (%)
0,1
0,12
Cj (at %)
Figure 7.19 The constant of the vacancy elimination reaction rate (K) as a function of the bulk boron concentration cv, normalized by the quenched-in vacancy concentration, N0 (Fraczkiewicz et al., 1998). Source: With kind permission of Elsevier. 8 7 6 5 4 3 2 1 0
0
24 h 3 months Time of annealing (h)
10
Figure 7.20 Effect of time of annealing at 400 C on the intergranular boron concentration. Fe40Al alloy, doped with 200 appm of boron (Fraczkiewicz et al., 1998). Source: With kind permission of Elsevier.
The effect of boron on B2 compounds
269
Figure 7.21 Dislocations network due to the thermal vacancies elimination in Fe40Al “pure” alloy: (A) quenchedand (B) annealed 400 C/24 h (Fraczkiewicz et al., 1998). Source: With kind permission of Elsevier.
Figure 7.22 Dislocations network due to the thermal vacancies elimination in the Fe40Al alloy, doped with 200 appm of B: (A) quenched and (B) annealed during 1 week at 400 C (Fraczkiewicz et al., 1998). Source: With kind permission of Elsevier.
as indicated by shrinkage measurements, the rearrangement of the structure on microscopic scale is slowed down. Disordering of the FeAl alloys is observed by addition of B, however, due to the low B concentration it was difficult to definitely decide if the entire FeAl matrix is disordered. At this stage it is suggested that B doped FeAl should contain some disordered B-rich regions allowing faster diffusion and so accelerating thermal vacancies elimination. The disordered region can be actually dislocation walls with high dislocation density.
270
Basic Compounds for Superalloys
References Crimp, M.A., Vedula, K., 1986. Mater. Sci. Eng. 78, 193. Deevi, S.C., Sikka, V.K., Inkson, B.J., Cahn, R.W., 1997. Scr. Mater. 36, 899. Fraczkiewicz, A., Gay, A.-S., Biscondi, M., 1998. Mater. Sci. Eng., A 258, 108. George, E.P., Liu, C.T., 1990. J. Mater. Res. 5, 754. Xie, C.Y., Jiang, B.H., Hu, G.X., 1993. Scr. Metall. Mater. 28, 1101.
Further Reading Paris, D., 1975. Scr. Metall. 9, 1373. Schulson, E.M., D.R. Barker, Scr. Metall. 17, 519 (1983).
7.1
7
NiAl
As indicated earlier ordered B2 NiAl has excellent properties among them low density, good oxidation and corrosion resistance and high melting point, therefore a possible high temperature structural material for industrial applications. However, the inherent brittleness at low temperature limited its use for aeroengine manufacturing. The grain boundaries were considered the weak region of NiAl where intergranular cohesion is not sufficiently strong and the failure by fracture on the application of load is considered to be intergranular. The addition of proper alloying elements is an effective method to improve the mechanical properties of intermetallic compounds such as NiAl. Such additives not only strengthen the matrix, but also significantly enhance the intergranular cohesion due to their segregation in grain boundaries. B is considered an excellent strengthener in various intermetallic compounds among them NiAl also. Experiments were performed to evaluate the effectiveness of B to suppress grain boundary fracture in NiAl and correlate it with grain boundary segregation and the improvement of the ductility. Table 7.1 shows the room temperature properties with and without B additions. C and Be are included in the table, since first principle calculations indicated that C is better than B in improving cohesive strength and Be was found to have a beneficial effect in the case of Ni3Al. One can observe in the table that the UTS is increased, both in the partially relaxed and relaxed stage. The expectation of improved ductility was not fulfilled when 300 ppm B was added, but grain boundary fracture was suppressed almost completely. Actually it is claimed that the B doped NiAl has a lower ductility than the undoped NiAl. In the experiments performed not showing improved ductility, can be explained by the fact that the high tensile strength in general is on the expense of elongation, namely the variation of strength is found to be opposite to the ductility. Increase in strength generally is not accompanied with increase in ductility. The addition of 300 ppm B, however, suppresses grain boundary fracture. Fig. 7.1 shows room-temperature fracture surface of B free and B containing NiAl. In order to determine if there was any change in the composition at grain boundaries relative to the bulk due to B additions plots of the peak height ratios (PHRs), namely that of Ni (848 eV)/Al (1396 eV) for both intergranular and transgranular areas was made for the alloys. This is seen in Fig. 7.2. Clearly the elemental peak heights were determined by Auger spectra. The Ni and Al peak spectra in the B doped NiAl at grain boundaries are higher than are without doping, an indication for B segregation in grain boundaries. No impurities Basic Compounds for Superalloys. DOI: https://doi.org/10.1016/B978-0-12-816133-3.00007-8 © 2018 Elsevier Inc. All rights reserved.
254
Basic Compounds for Superalloys
Room-temperature tensile properties of NiAl microalloyed with B, C, and Be (George and Liu, 1990)
Table 7.1
Alloy composition (at.%)
Microstructure (heat treatment)
Yield strength (MPa)
Ultimate tensile strength (MPa)
Tensile elongationa (%)
Ni50Al Ni50Al NiAl 1 300 ppm Bb NiAl 1 300 ppm Bb NiAl 1 300 ppm Cb NiAl 1 300 ppm Cb NiAl 1 500 ppm Beb NiAl 1 500 ppm Beb
Partial rex. (1 h/500 C) Rex. (1 h/805 C) Partial rex. (1 h/800 C) Rex. (1 h/1000 C) Partial rex. (1 h/500 C) Rex. (1 h/1100 C) Partial rex. (1 h/950 C) Rex. (1 h/1050 C)
198 154
262 229 400 329 429 336 268 307
2.0 2.2
c c c c
188 178
c c c c
1.9 3.0
a
Elongation measured from a strip chart. Weight parts per million. Fracture prior to macroscopic yielding. Source: With kind permission of Cambridge University Press.
b c
Figure 7.1 Room-temperature fracture surfaces investigated in this study (George and Liu, 1990). Source: With kind permission of Cambridge University Press.
were detected in the stoichiometric NiAl which could possibly explain the intergranular fracture, an indication that NiAl is intrinsically brittle. Of interest is to note that the Ni-rich B2 NiAl shows not only martensitic transformation, but also perfect pseudoelasticity. The 37.5 at.% Al alloy is off-stoichiometry, but still having a B2 structure. B is added because it is one of the most effective microalloying elements improving the ductility and suppressing the intergranular fracture. It changes the Ms temperature but to a smaller degree the Mf temperature and influences the reverse transformation temperature. Also the B addition seems to restrict grain growth to some extent and the 37.5 at.% AlB has a smaller grain size. X-ray diffraction indicates that both alloys are ordered B2 structures. The phase transitions was evaluated from expansion vs temperature curves, shown for the 37.5 at.% Al and the 37.5 at.% AlB NiAl alloys in Fig. 7.3.
The effect of boron on B2 compounds
255
NiAl
NiAl + 300 ppm B 4.0 IG TG
3.5
Peak height ratio (Ni/Al)
Peak height ratio (Ni/Al)
4.0
3.0 2.5 2.0 1.5 1.0
3.0 2.5 2.0 1.5 1.0
0
5
10
15
20
25
IG TG
3.5
0
10
Analysis No.
20
30
Analysis No.
Figure 7.2 Ni/Al PHRs for various intergranular and transgranular areas in the alloys analyzed by Auger electron spectroscopy (George and Liu, 1990). Source: With kind permission of Cambridge University Press.
(A)
As Linear expansion
Mf Af Ms
(B) Mf.As
Af Ms –200
–100 0 Temperature (°C)
100
Figure 7.3 (A) Linear expansion-temperature curves of the Ni-37.5 at.% Al specimen and (B) the Ni-37.5 at.% AlB specimen (Xie et al., 1993). Source: With kind permission of Elsevier.
Note that B additions modify the Ms and the Af temperatures, but almost no change has occurred in the Mf and As temperatures (Table 7.2). The As refer to the reverse transition (in FeC phase diagram, it refers to Austenite). Fracture occurs readily by a very small force without B addition due to the extreme brittleness of the annealed Ni 37.5 at.% alloy. The intergranular fracture of the
256
Basic Compounds for Superalloys
Phase transformation temperatures of the alloys (Xie et al., 1993)
Table 7.2
Alloy
Ms ( C)
Mf( C)
As ( C)
Af( C)
Ni37.5 at.% Al Ni37.5 at.% AlB
2 118 2 84
2 129 2 129
2 121 2 129
2 111 2 75
Source: With kind permission of Elsevier.
Figure 7.4 (A) Morphologies of the fracture surface of the Ni-37.5 at.% Al specimen and (B) the Ni-37.5 at.% AlB specimen (Xie et al., 1993). Source: With kind permission of Elsevier. 700
Stress (MPa)
600 (A)
500
(B)
400 300 200 100
0.25 % 0
0 Strain
Figure 7.5 Compressive stressstrain curves of the Ni-37.5 at.% AlB specimens (Xie et al., 1993). Source: With kind permission of Elsevier.
unalloyed Ni 37.5 at.% Al is seen in Fig. 7.4A and some ductility was observed in the 37.5 at.% AlB NiAl, shown for comparison in Fig. 7.4B. The transformation temperatures of the alloys marked in Fig. 7.3 are shown in Table 7.2. Compressive stressstrain curves of the B doped alloy are illustrated in Fig. 7.5.
The effect of boron on B2 compounds
257
There is a definite yield point at 0.9% strain, as marked by the arrow and a total strain of 3.2% is obtained before fracture. The strain is recovered completely on unloading of the specimen compressed to 2.8% strain as seen in Fig. 7.5B. This observation of perfect pseudoelasticity is similar to the observations in shape memory alloys. Thus a stress induced martensitic transformation is taking place during loading and its reverse deformation occurs on unloading. B modifies the Ms temperature and induces a complete pseudoelasticity strain of 2.8% at room temperature.
7.2
FeAl
As in other B2 compounds additions of small amounts of B may change the fracture mode from intergranular to transgranular and induce ductility in these alloys at room temperature. Room temperature ductility is clearly very important, allowing machinability in inherently brittle materials, by alloying with small amounts of an additive. In the following information on FeAl alloys of different Al and B concentrations is presented. Due to the large homogeneity range showing an extended phase field in FeAl at the location of B2, it is possible to obtain iron aluminides of off-stoichiometric compositions, where the deficiency of atomic species leads to constitutional vacancies formation in order to maintain charge neutrality. It is known that thermal vacancies harden aluminides and quenched in vacancies can be relieved by annealing at some temperature. The concentration of vacancies was reported to increase with Al content. The hardening of FeAl depends on the substitutional element. In this book the concern is the effect of B as the additional element. The reports on the effect of B on the strengthening of FeAl are not uniform and these are considered in the following text. B in the amounts of 0.05, 0.1, 0.15, 0.3, 0.5, 1, 2, and 4 at.% were added to the base Fe45 at.% Al alloy. As cast Rockwell hardnesses of FeAl with B additions were measured prior to heat treatment of each specimen at 950 C for 2 h in air followed by a vacancy annealing of 400 C for 120 h in air. Hardnesses were measured also after vacancy annealing. The effect of B on the Rockwell hardnesses is shown in Fig. 7.6 indicating the as-cast and heat treated hardnesses. The hardness of the as cast specimens is the highest at all B concentrations. However, the second anneal at 950 C for 2 h and 400 C for 120 h increased the hardnesses at all B concentration but to a level still lower than the as cast specimens. Also note that the hardness of B containing alloys at all B levels is lower than the unalloyed FeAl. It is interesting to note that the 0.2% yield strength is significantly higher than the unalloyed FeAl as seen in Fig. 7.7. B doped alloys show only a very small decrease in yield strength after the vacancy-annealing treatment. The oil-dipped specimens indicated in Fig. 7.7 was intended to eliminate possible hydrogen embrittlement and notch effects to which FeAl tested in air is very sensitive. No meaningful difference in strength between the oil dipped and undipped specimens is observed. The effect of B on tensile elongation is seen in Fig. 7.8. The tensile elongation is low in both, the
258
Basic Compounds for Superalloys
80 As Cast 1st-950°C/2 h 1st-400°C/120 h 2nd-950°C/2 h 2nd-400°C/120 h
Hardness (Rc)
60
40
20
0
0
0.05
0.1 0.15 % Boron
0.3
0.5
Figure 7.6 Effect of boron content on the hardnesses of Fe45 at.% Al alloys in the as-cast condition and after annealing at 950 C for 2 h, then air cooling, and subsequent aging at 400 C for 120 h (Deevi et al., 1997).
0.2% Yield strength (MPa)
750 950°C/2 h Air cool 400°C/120 h Air cool 950°C/2 h Air cool-Oil Dip 400°C/120 h Air cool-Oil Dip
500
250
0 Fe–-Al
0.05 0.1 Boron content (at.%)
0.15
Figure 7.7 Effect of boron content on room-temperature yield strengths of Fe45 at.% Al alloys (Deevi et al., 1997).
B-doped and B-free specimens, but it increases after the vacancy annealing treatment (at 400 C for 120 h) to 0.46%. The alloys with 0.1 at.% B had the highest elongation of 1.9% after vacancy annealing treatment. The microstructures of annealed and vacancy annealed treated samples were quite similar, but a small grain size reduction was observed in the B-doped specimens in the range of 0.050.15 at.% B. The increase in room temperature properties such as yield strength and UTS by B addition is attributed to the grain boundary strengthening (Table 7.3).
The effect of boron on B2 compounds
259
Total elongation (%)
3 950°C/2 h Air cool 400°C/120 h Air cool 950°C/2 h Air cool-Oil Dip 400°C/120 h Air cool-Oil Dip 2
1
0
0
0.05 0.1 Boron content (at.%)
0.15
Figure 7.8 Effect of boron content on the room-temperature tensile elongations of Fe45 at.% Al alloys (Deevi et al., 1997).
B is claimed to segregate to grain boundaries which is considered by some to be the cause of the strengthening effect. The room temperature fracture mode changes at high temperatures from intergranular to transgranular, namely from brittle to ductile, as observed in both alloys the alloy of Fe46.1 at.% Al and the Fe46.1 at.% Al-0.78 at.% B. The engineering stressstrain curves for these two alloys at different temperatures are shown in Fig. 7.9. The tensile properties are listed in Table 7.2. The grains were equiaxed after the homogenization treatment in both alloys indicating recrystallization as seen in the microstructure of Fig. 7.10. The second-phase particles in the FeAlB specimens are suspected to be borides. In the FeAl specimens—which were totally brittle up to 600K—a significant ductility of 8% elongation in one specimen and 55% plastic elongation to fracture in another were measured. Disappointedly the FeAlB alloy remains brittle at 300K, but shows some ductility at 560K. Thus only a modest increase in ductility is the result of B addition. However a significant increase in the strength of FeAl is obtained by the addition of B. Also the addition of B changes the fracture mode from intergranular to transgranular, but the alloy remains brittle. The fracture surfaces of these two alloys are shown in Fig. 7.11 and as seen, at 300K, the FeAl specimen is totally intergranular, whereas the FeAlB specimen displays totally transgranular cleavage fracture. At 640K—just above the ductile-to-brittle transition temperature—the intergranular mode for FeAl and transgranular mode for FeAlB is maintained and still observed but some ductility is present. The FeAl specimen which showed a large ductility at 640K is characterized by dimple rupture features as seen in Fig. 7.12. The dislocation structures of homogenized FeAl and FeAlB specimens are illustrated in Fig. 7.13. The dislocations were introduced during the hot extrusion process and were sessile in their nature to be removed by heat treatment. They had
Results of tensile tests on FeAl and FeAlB as a function of temperature (constant strain rate, about 3 3 1023s21) (Crimp and Vedula, 1986)
Table 7.3
Temperature (K)
FeAl Yield strengtha (klbf in22)
300 560 600 640 (1) 640 (2) a
Ultimate tensile strengthb (klbf in22)
Fracture stressc (klbf in22)
FeAlB Elongationd (%)
Yield strengtha (klbf in22)
Ultimate tensile strengthb (klbf in22)
77 80 89 71
92 85
8 55
137 120 116
Fracture stressc (klbf in22) 141 143 130
129
Elongationd (%)
,1 2.5 11
0.2% offset yield strength for specimens which exhibited plastic yielding. Ultimate tensile strength for specimens which showed maximum load followed by necking elongation. Stress at which brittle fracture occurred before yielding for brittle specimens or stress at which ductile fracture occurred before necking elongation for ductile specimens. d Percentage elongation to failure in a gauge length of 3 cm for ductile specimens. Source: With kind permission of Elsevier. b c
The effect of boron on B2 compounds
261
–2
klbf in
MPa
160 1000
Stress
120 640 k 600 k 300 k
80
640 k 0.55 Strain at 500 fracture
40
0 0
0.04
0.08
(A)
0.12 Strain
0.16
0.20
klbf in–2
MPa
160 300 k 560 k 600 k
1000 640 k
Stress
120
80
500
40
0 (B)
0.04
0.08
0.12 Strain
0.16
0.20
Figure 7.9 Comparison of engineering stressstrain curves for (A) FeAl and (B) FeAlB alloy (Crimp and Vedula, 1986). Source: With kind permission of Elsevier.
definite orientations and the dislocations consisted of sharp jogs without evidence of faults or superlattice dislocations. Addition of B had a noticeable effect on the substructure of the dislocations by decreasing the tendency of having a specific orientation, making the jogs smoother and in inducing stacking faults formation. Because of B segregation in grain boundaries a strengthening of boundaries relative to the matrix was observed, as a consequence of modifying the fracture mode from intergranular to transgranular. However, no measurable ductility at room
262
Basic Compounds for Superalloys
Figure 7.10 Comparison of optical micrographs of the two alloys, illustrating recrystallized grains and prior particle boundaries in both alloys: (A) FeAl, longitudinal, (B) FeAl, transverse, (C) FeAlB, longitudinal, and (D) FeAlB, transverse (Crimp and Vedula, 1986). Source: With kind permission of Elsevier.
temperature was observed despite the change occurring in the fracture mode. Despite the relatively high level of B addition (0.78 at.%) which is beyond the amount required for strengthening grain boundary cohesion, it is speculated that cross-slip may contribute to early crack formation before plastic deformation could take place across grain boundaries. The concept is thus associated with the stress required to initiate plastic flow. There is a competition between this stress and the one required for brittle fracture either intergranular or transgranular. If the stress required to initiate brittle fracture (intergranular or transgranular) is less than that required to cause plastic flow, the specimen will break in a brittle manner, whereas in the opposite case (stress for plastic flow lower than the brittle fracture stress) as it is at higher temperatures the specimen exhibits ductility. The addition of B modifies the dislocation structure by forming stacking faults, which without B additions is a material without faulted structure. The presence of the stacking faults makes slip more difficult resulting in the increase in the matrix strength. The possibility of solid solution strengthening by B and particle strengthening are not ruled out. Thus
The effect of boron on B2 compounds
263
Figure 7.11 Comparison of scanning electron micrographs of the fracture surfaces for the two alloys at two different temperatures: (A) FeAl, 300K; (B) FeAl, 640K; (C) FeAlB, 300K; (D) FeAlB, 640K (Crimp and Vedula, 1986). Source: With kind permission of Elsevier.
a twofold strengthening by B additions should be noted: (a) B segregation at grain boundaries and modifying at low temperatures the fracture mode from intergranular to transgranular and (b) matrix strengthening by the formation of stacking faults making cross-slip more difficult. A different interpretation of the effect of B addition involves the influence of vacancies in the FeAl. Large grain alloys (B400 μm) containing 40, 46, or 50 at.% Al with additions of 80200 appm (atomic parts per million) are considered. A two-stage heat treatment was applied to these alloys: (a) to induce thermal vacancies at 950 C/1 h to obtain a maximum concentration of vacancies followed by air quenching and (b) 400 C/24 h for the elimination of excess thermal vacancies. Experimental results showing B concentration was evaluated either by the McLean or the Fowler’s model given as McLean
ΔGsegr cj cv 2 cj 5 exp 2 1 2 cv kT c0j
(7.1)
264
Basic Compounds for Superalloys
Figure 7.12 Scanning electron microscopy fracture surface of FeAl alloys showing (A) large ductility and (B) typical ductile fracture at 640K (Crimp and Vedula, 1986). Source: With kind permission of Elsevier.
Fowler
2 ΔGsegr 2 zw:cj =c0j cj cv 2 c 5 exp j 1 2 cv kT c0j
! (7.2)
where cj is the intergranular segregation of the segregating atoms,c0j is the saturation level at grain boundaries (considered to be a monolayer), cv is the bulk concentration representing the solubility limit, ΔGsegr is the energy of segregation and zw is an interaction factor. According to the sign of the zw parameter, the segregation may be limited (zw . 0-repulsive interactions) or enhanced (zw , 0-attractive interactions). The effect of bulk concentration, cv on intergranular segregation, cj is illustrated in Fig. 7.14. The intergranular concentration cj as a function of temperature is seen in Fig. 7.15. No significant influence by heat treatment in the temperature range 300950 C could be observed and the segregation level is almost the same as in the quenched alloy. This observation does not support an equilibrium segregation model. Both the equilibrium model of McLean and Fowler predict a strong decrease of the intergranular model with temperature [Eqs. (7.1)and (7.2)].
The effect of boron on B2 compounds
265
Figure 7.13 Transmission electron micrographs illustrating the major differences in the substructure of the two alloys (A) FeAl and (B) FeAlB (Crimp and Vedula, 1986). Source: With kind permission of Elsevier.
Because of the large difference in the energies of formation and migration, (0.7 and 1.7 eV/atom, respectively) the quenched alloys are able to retain the high concentration of thermal vacancies up to 0.2 at.% in Fe40Al and even up to 2 at.% in Fe50Al alloys (Paris). These vacancies—in excess at room temperature—may be eliminated at a low temperature heat treatment. The elimination of the vacancies can be evaluated from density measurements. This is shown for pure and B doped FeAl of 40, 46, and 50 at.% Al in Figs. 7.167.18, respectively. As indicated earlier the alloys were air quenched from 950 C. Under the assumption that one controlling mechanism of vacancy elimination is dominant, the decrease in vacancy concentration dN/dN0 may be expressed as dN K 5 2 exp dN0 kT
(7.3)
266
Basic Compounds for Superalloys
18 Exp (max) Fowler McLean
16 14 Cj (at%)
12 10 8 6 4 2 0 500
0
1000 Cv (appm)
1500
2000
Figure 7.14 Effect of the bulk boron concentration (cv) on its intergranular concentration (cj) in Fe40Al alloys. Heat treatment: 950 C/1 h air quench 1 400 C/24 h (Fraczkiewicz et al., 1998). Source: With kind permission of Elsevier. 9 8 7 Cj (at%)
6 5 4 3 2 1 0 200
300
400
500 600 700 Temperature (°C)
800
900
1000
Figure 7.15 Effect of temperature on the intergranular concentration of boron. Fe45Al 1 400 appm B alloy; annealing during 24 h (Fraczkiewicz et al., 1998). Source: With kind permission of Elsevier.
K is the reaction constant. The ratio of Kdoped/Kpure as a function of bulk B concentration is given in Fig. 7.19. Clearly the time of annealing influences the concentration of the vacancies as can be seen in Fig. 7.20 when it is carried out at 400 C. The acceleration of vacancy elimination kinetics in the presence of boron can be explained by the formation of boron-vacancy complexes, migrating to the grain boundaries faster than the desegregated B returning back to the matrix. Nevertheless, still an equilibrium segregation of B at a low level remained. A new dislocation structure is formed following the elimination of thermal vacancies. In the pure FeAl, vacancy elimination following quenching, is associated already at the start with dislocation loops creation with h111i Burgers vector. But
The effect of boron on B2 compounds
267
Atomic fraction of vacancies (%)
0,25 Fe–40Al–200B Fe–40Al–40B
0,2
Fe–40Al
0,15 0,1 0,05 0 0
2
4
6 8 10 Annealing time (h)
12
14
16
Figure 7.16 Kinetics of quenched-in vacancies elimination during an isothermal annealing at 380 C. Dilatometric study of Fe40Al alloys, pure or B-doped (Fraczkiewicz et al., 1998). Source: With kind permission of Elsevier.
Atomic fraction of vacancies (%)
0,45 Fe–46Al–200B Fe–46Al
0,4 0,35 0,3 0,25 0,2 0,15 0,1 0,05 0
0
5
10
15 20 Annealing time (h)
25
30
35
Figure 7.17 Kinetics of quenched-in vacancies elimination during isothermal annealing at 380 C. Dilatometric study of Fe46Al alloys, pure or B-doped (Fraczkiewicz et al., 1998). Source: With kind permission of Elsevier.
after annealing the homogeneous h100i dislocation network is obtained. This can be seen in Fig. 7.21. In B doped FeAl the dislocation structure is different as shown in Fig. 7.22. Vacancy elimination starts faster than in pure FeAl. As a result the dislocation network due to vacancies elimination is almost the same in the quenched state as in the annealed samples for 24 h at 400 C (Fig. 7.22A); it is regular but heterogeneous. The cell walls are B3 μm in size with only a few individual dislocations (namely high dislocation density walls). After 400 C annealing for a week a homogeneous structure similar to the pure alloy is obtained as seen in Fig. 7.22B. Despite the acceleration macroscopic effect of vacancy elimination by B addition
268
Basic Compounds for Superalloys
Atomic fraction of vacancies (%)
0,8 Fe–50Al–200B Fe–50Al Fe–50Al–800B
0,7 0,6 0,5 0,4 0,3 0,2 0,1 0 0
10
20 30 40 Annealing time (h)
50
60
Kdoped/Kpure
Figure 7.18 Kinetics of quenched-in vacancies elimination during isothermal annealing at 380 C. Dilatometric study of Fe50Al alloys, pure or B-doped (Fraczkiewicz et al., 1998). Source: With kind permission of Elsevier. 4,5 4 3,5 3 2,5 2 1,5 1 0,5 0
Fe–40AI Fe–46AI Fe–50AI
0
0,02
0,04
0,06 0,08 Cv/No (%)
0,1
0,12
Cj (at %)
Figure 7.19 The constant of the vacancy elimination reaction rate (K) as a function of the bulk boron concentration cv, normalized by the quenched-in vacancy concentration, N0 (Fraczkiewicz et al., 1998). Source: With kind permission of Elsevier. 8 7 6 5 4 3 2 1 0
0
24 h 3 months Time of annealing (h)
10
Figure 7.20 Effect of time of annealing at 400 C on the intergranular boron concentration. Fe40Al alloy, doped with 200 appm of boron (Fraczkiewicz et al., 1998). Source: With kind permission of Elsevier.
The effect of boron on B2 compounds
269
Figure 7.21 Dislocations network due to the thermal vacancies elimination in Fe40Al “pure” alloy: (A) quenchedand (B) annealed 400 C/24 h (Fraczkiewicz et al., 1998). Source: With kind permission of Elsevier.
Figure 7.22 Dislocations network due to the thermal vacancies elimination in the Fe40Al alloy, doped with 200 appm of B: (A) quenched and (B) annealed during 1 week at 400 C (Fraczkiewicz et al., 1998). Source: With kind permission of Elsevier.
as indicated by shrinkage measurements, the rearrangement of the structure on microscopic scale is slowed down. Disordering of the FeAl alloys is observed by addition of B, however, due to the low B concentration it was difficult to definitely decide if the entire FeAl matrix is disordered. At this stage it is suggested that B doped FeAl should contain some disordered B-rich regions allowing faster diffusion and so accelerating thermal vacancies elimination. The disordered region can be actually dislocation walls with high dislocation density.
270
Basic Compounds for Superalloys
References Crimp, M.A., Vedula, K., 1986. Mater. Sci. Eng. 78, 193. Deevi, S.C., Sikka, V.K., Inkson, B.J., Cahn, R.W., 1997. Scr. Mater. 36, 899. Fraczkiewicz, A., Gay, A.-S., Biscondi, M., 1998. Mater. Sci. Eng., A 258, 108. George, E.P., Liu, C.T., 1990. J. Mater. Res. 5, 754. Xie, C.Y., Jiang, B.H., Hu, G.X., 1993. Scr. Metall. Mater. 28, 1101.
Further Reading Paris, D., 1975. Scr. Metall. 9, 1373. Schulson, E.M., D.R. Barker, Scr. Metall. 17, 519 (1983).