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Effect of Thermomechanical Treatments on the Room- Temperature Mechanical Behavior of Iron Aluminide Fe3AI

ARVIND AGARWAL, R. BALASUBRAMANIAM, and S. BHARGAVA

The room-temperature hydrogen embrittlement (HE) problem in iron aluminides has restricted their use as high-temperature structural materials. The role of thermomechanical treatments (TMT), i.e., rolling at 500 ~ 800 ~ and 1000 ~ and post-TMT heat treatments, i.e., recrystallization at 750 ~ and ordering at 500 ~ in affecting the room-temperature mechanical properties of Fe-25A1 intermetallic alloy has been studied from a processing-structure-properties correlation viewpoint. It was found that when this alloy is rolled at higher temperature, it exhibits a higher fracture strength. This has been attributed to fine subgrain size (28 /x) due to dynamic recrystallization occurring at the higher rolling temperature of 1000 ~ However, when this alloy is rolled at 1000 ~ and then recrystallized, it shows the highest ductility but poor fracture strength. This behavior has been ascribed to the partially recrystallized microstructure, which prevents hydrogen ingress through grain boundaries and minimizes hydrogen embrittlement. When the alloy is rolled at 1000 ~ and then ordered at 500 ~ for 100 hours, it shows the highest fracture strength, due to its finer grain size. The alloy rolled at 500 ~ and then ordered undergoes grain growth. Hence, it exhibits a lower fracture strength of 360 MPa. Fracture morphologies of the alloy were found to be typical of brittle fracture, i.e., cleavage-type fracture in all the cases.

I. INTRODUCTION

ORDERED iron aluminide intermetallics of composition Fe3A1 and FeAI possess attractive properties for appli- cation as structural materials at elevated temperahares in aggressive environments.I ~z3~ However, their poor room- temperature ductility limits their use as engineering materials as they are difficult to process into useful shapes such as plates and tubes. In recent years, efforts have been in- tensified to identify both the intrinsic as well as the extrinsic factors governing their room-temperature brittle fracture. [41

Several studies have established that the room-temperature ductility is caused mainly by an extrinsic effect--en- vironmental embrittlement due to hydrogen. Is,6.n Several methods have been proposed to minimize hydrogen embrittlement (HE) in iron aluminides. Most of the methods that have been suggested to curb HE aim to restrict entry of hydrogen into the lattice by providing a passive film on the iron aluminide surface. Oxide coatings have been found to be beneficial in increasing the ductility of iron aluminides.t81 Another method for improving the ductility of iron aluminides is by the addition of chromium. Even small chromium additions are effective in minimizing HE and providing ductility.tgI The role of chromium in affecting HE and causing ductility enhancement has been elucidated from an electrochemical viewpoint.rio] It has been recently shown that thermomechanical treatments (TMT) also play a crucial role in affecting the mechanical properties of Fe- 25A1-1 B intermetallic alloy, vq

The present investigation aims at understanding the ef-

ARVIND AGARWAL, Graduate Student, R. BALASUBRAMANIAM, Assistant Professor, and S. BHARGAVA, Professor, are with the Department of Materials and Metallurgical Engineering, Indian Institute of Technology, Kanpur 208 016, India.

Manuscript submitted February 21, 1995.

fect of various TMT on the mechanical behavior of Fe-25AI iron aluminide.

II. EXPERIMENTAL PROCEDURE

The starting intermetallic alloy was of composition Fe- 25A1 and was supplied by the Defence Metallurgical Re- search Laboratory (DMRL) (Hyderabad) in the form of a cast pancake of about 120-mm diameter and 15-mm height. The composition of the alloy was analyzed by a JEOL*

*JEOL is a trademark of Japan Electron Optics Ltd., Tokyo.

electron probe microanalyzer and was found to be 74.72Fe- 25.28A1 (_+ 0.25). Rectangular pieces of 10-mm thickness were cut from this pancake and subsequently homogenized at 1000 ~ for 4 hours prior to their rolling via different thermomechanical schedules. Various thermomechanical schedules followed in the present study are schematically shown in Figure 1. These schedules basically consisted of deformation of the Fe-25AI alloy in three different phase fields, namely, (1) in the disordered a phase field at 1000 ~ (2) in the ordered B2 phase field at 800 ~ and (3) in the ordered DO 3 phase field at 500 ~ Prior to rolling, the samples were soaked in a furnace in an air atmosphere at their respective rolling temperatures. The desirable thickness deformation (80 pct) to each sample was achieved by multipass rolling. In order to remove the thin surface scale formed during the course of hot rolling, the specimens were polished on emery paper prior to their post-TMT treatments and tensile testing. Post-TMT treatments consisted of (1) recrystallization annealing at 750 ~ for 1 hour and (2) ordering treatment at 500 ~ for 100 hours. The specimens were encapsulated in quartz tube under vacuum for all post- TMT heat treatments.

X-ray diffraction (XRD) patterns of the Fe-25A1 alloy were recorded using Cu K~ radiation on a Rich-Siefert

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 27A, OCTOBER 1996---2985

L PIates of Fe-25At J

] Homogen iza t ion a t 1000~ for 4 hours

Rolting at 1000~ in ol phase field

l Rolling cat 800~ in B2 phase field

Rolling at 500~ n DO3 pho, se field

l Recrystallization at 750~ for 1hour [

Ordering at 500~ I for 100 hours [

I i

CharacterizationJ

Mechonical testing (and fractogrophy

Fig. 1--Flow chart of the experimental procedure adopted in the present study.

s~

c-

(220)

C400)

Rolled + ordered (422) I /

C620) Y ~ k

Rolled + tee

As roiled

A s r e c e t v e d

136 134 "118 116 114 "100 98 96 44 42 40 38 2 e ( d e o r ~ )

Fig. 2--XRD patterns obtained from the Fe-25A1 alloy: in the as-received condition, after rolling at 1000 ~ to 80 pct thickness reduction, after rolling at 1000 ~ to 80 pct tluckness reduction + recrystallizing at 750 ~ for 1 h, and after rolhng at 1000 ~ to 80 pct thickness reduction + ordering at 500 ~ for 100 h.

2002 X-ray diffractometer. Specimens used for this purpose were in the form of strips, before and after post-TMT heat treatments. Microhardness of these specimens was also measured using a Carl Zeiss Jena optical microscope at- tached with a microhardness tester. In all the cases, inden-

Table I. Microhardness of Fe-25AI Alloy after Various Thermomechanical and Heat Treatments

Stage of Processing Microhardness (VHN)

As-received 225 ___ 8 Rolled* 390 ___ 14

Rolled + recrystallized 345 + 10 Rolled + ordered 380 _+ 10

*The rolling temperatures were not found to influence the microhardness of the Fe-25A1 alloy.

tations were made using a load of 100 g and hardness values obtained as Vickers hardness numbers (VHN).

Tensile-test specimens of size 10 by 5 by 1.25 mm were machined from the strips. In order to remove the thin surface scale formed during rolling and the surface flaws in- troduced during machining, the tensile-test specimens were carefully polished on emery paper prior to their testing. Tensile tests were carried out in air at room temperature on an Instron machine at a constant strain rate of 10 4 s-1. Moreover, in order to minimize the entry of hydrogen into the lattice during straining, the specimens were coated with silicon oil prior to their testing. Because of the shortage of the thermomechanically treated alloy, larger numbers of tensile tests to generate the scatter in the data could not be undertaken. Duplicate specimens from the as-rolled strips and one specimen from the heat-treated strips were tested.

Microstructures of samples after different stages of their processing were observed using a Leitz Laborlux microscope. Metallography samples were prepared by standard methods. Etching response of samples, however, was found to be somewhat complicated and grain boundaries could be best revealed by the etchant, which consisted of eight parts of CH3COOH, four parts of HNO3, and one part of HC15 TM

Optimum results were obtained by keeping the etching time between 5 and 15 seconds.

Fractured tensile-test specimens were examined under a JEOL JSM 840A scanning electron microscope (SEM). In order to study the nature of the cleavage planes, the fracture surfaces were studied in etched and unetched conditions. Etch- ing of the previously observed fracture surface was carried out using the aforementioned etchant for 1 minute. The etch pits were observed under the SEM to reveal the crystallographic nature of the cleavage planes on the fracture surface.

III. RESULTS AND DISCUSSION

A. X-Ray Diffraction Analys&

The XRD pattems obtained from the Fe-25A1 alloy at different stages of its processing are shown in Figure 2. It can be seen that in the as-received condition the alloy showed peaks corresponding to fundamental reflections such as (220), (400), (422), (440), (620), and (444). Oki et aL [~31 have shown that the extinction rules to be followed are as follows: h, k, and l are all even and (h + k + /)/2 is even for the disordered a phase, h, k, and l are all even and (h + k + /)/2 is odd for the imperfectly ordered B2 phase, and h, k, and l are all odd for the DO 3 phase. It was

2986--VOLUME 27A, OCTOBER 1996 METALLURGICAL AND MATERIALS TRANSACTIONS A

Ul ul

t - " 0 L o t- o u

o ~ ~E

36q

3 5 0

344

3 3 0

32(;

31q

30 ( f ] I t 0 0.50 1.00 1.50 ZOO 2.50

Annealing Time (hr) (a)

(b)

(c)

(a)

Fig. 3-- (a) Microhardness of the Fe-25A1 alloy as a function of annealing time after rolling at 500 ~ and recrystallizing at 750 ~ Optical micrographs of the alloy annealed for (b) 0.5 h, (c) 1 h, and (d) 2 h at 750 ~

therefore concluded that in the as-received condition, the alloy essentially consisted of disordered bcc structure.

The XRD patterns from the alloy after its TMT at three different rolling temperatures and post-TMT recrystallization anneal at 750 ~ did not show any significant variation. The pattems obtained from the specimen rolled at 1000 ~ and from the specimen rolled at 1000 ~ and recrystallized at 750 ~ are shown in Figure 2. It can be seen that the alloy after its hot rolling and recrystallization anneal had a phase distribution similar to that of the as-received material.

The ordering treatment in Fe3Al-base alloys has earlier been investigated by several authors, t~3 ~71 From changes occurring in electrical resistivity during isothermal heat-treat-

ment experiments, McQueen and KuczynskitlS] concluded that transformation from the disordered phase or the B2 phase to the D O 3 phase occurred by a nucleation and growth process, and their results were substantiated by Da- vies tlr] and SelisskiyY] Oki et a/. [13,14] studied the transformation behavior of Fe3Al-base alloy by XRD and concluded that (1) the disordered phase could not be re- tained even by quenching the alloy from the disordered phase field and small domains of B2 invariably formed, (2) transformation from the B2 structure to the D O 3 structure during isothermal treatments occurred by the nucleation and growth of DO3 domains in the B2 matrix, (3) the superlattice reflections from the DO3 phase were practically too

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 27A, OCTOBER 1996--2987

weak except for the (111) reflection in some cases, and (4) broadening of the fundamental peak corresponding to (440) reflection occurred upon ordering. The broadening of the (440) reflection was also observed by Selisskiy. v51 Oki et al. suggested that the width of the (440) reflection was orig- inally small but increased rapidly and reached a maximum value as the alloy was annealed isothermally. On the basis of the experimentally observed difference in lattice para- meters of the B2 and the DO3 phases and the fact that the amount of change in the (440) line width coincided with the difference between the (440) peak positions for the B2 and the DO 3 structures, they concluded that the process of ordering was heterogeneous.

The XRD pattem obtained after the ordering treatment is shown in Figure 2. The DO3 superlattice reflections such as (111), (311), (331), etc., could not be observed. A possible reason for their absence could be that these peaks are of very low intensity and can be observed only after a high degree of ordering is achieved. However, broadening of the peak corresponding to (440) reflection can clearly be seen in the pattern and, as explained in the previous paragraph, can be taken as evidence for the coexistence of the DO 3 and the B2 phases after the ordering treatment performed on a thermomechanically treated Fe-25A1 alloy. It should be also be noted from Figure 2 that the broadening of the (440) peak did not occur in the rolled and recrystallized alloy. The ordering treatment at 500 ~ also showed an extra peak at 20 = 41.4 deg. This peak does not correspond to any reflecting plane in the cubic system and has not been indexed in Figure 2. The nature of the XRD pattems was also very similar for the alloy rolled at three different temperatures and then ordered.

B. Microhardness Measurements

Table I shows the microhardness values of the Fe-25A1 alloy at different stages of processing. The scatter in the microhardness data has been indicated in the table. As shown in the table, the alloy was relatively soft in the as- received cast condition (225 VHN). However, its microhardness increased substantially to 390 VHN after rolling at 1000 ~ Further, the rolling temperature did not affect the hardness of the alloy in any significant manner, as the microhardness values obtained for specimens rolled at 800 ~ and 500 ~ were similar.

The increase in microhardness can be explained by the high dislocation density created as a result of the deformation processing and fast cooling rates employed after hot rollingY 81 The rolled alloy retains a high vacancy concentration after quenching from high temperatures. Baker and Nagpal have shown that these quenched-in vacancies inter- act with dislocations and form dislocation loops, which result in increased hardness, and that the water-quenched iron aluminides possessed the highest hardness because of the fast cooling rate employed.t~81 Thus, it can be concluded that water quenching and the heavy deformation of 80-pct thickness reduction were responsible for the increase in the hardness of the rolled alloy compared to the as-received alloy. The microhardness of the rolled and recrystallized alloy is lower than that of the as-rolled alloy. This is expected, as there is a reduction in the dislocation density because of annihilation of dislocations and vacancies during annealing.

The recrystallization behavior of Fe3Al-base alloys is

Table II. Room-Temperature Tensile Properties of Fe-25AI Alloy after Various Thermomechanical and Heat Treatments

Ultimate Yield Tensile

Strength Strength Ductility Grain Condition ( M P a ) (MPa) (Pct) Fracture Size (p~)

R1000 830 840 0 cleavage 28 R1000 795 810 0 cleavage 28 R800 625 730 0 cleavage * R800 610 710 0 cleavage * R500 370 370 0 cleavage * R500 350 350 0 cleavage * R1000 + rec. 380 460 1.4 cleavage 295 R800 + rec. 580 570 0 cleavage 140 R500 + rec. 670 730 0.9 cleavage 90 R1000 + ord. 565 560 0 cleavage 155 R800 + ord. 455 455 0 cleavage 340 R500 + ord. 360 360 0 cleavage 720

*Grain size measurements were not possible in these cases because of the highly deformed and elongated structure.

greatly influenced by variables related to microstructures developed by TMTs, alloying, phases present, their volume fractions, and their defect structure.t~91 Thus, it is very difficult to predict the recrystallization response and kinetics in Fe3Al-base alloys at a given temperature. Therefore, in order to establish the recrystallization schedule for the purpose of mechanical testing, the samples rolled at 500 ~ were annealed at 750 ~ for 0.5, l, and 2 hours. The variation of microhardness of these recrystallized samples with annealing time is presented in Figure 3(a). It is clear that the hardness of the alloy decreased with increasing recrystallization time at 750 ~ The variation in the hardness of the alloy can be understood in terms of its microstructures, which were obtained after different annealing times and are presented in Figures 3(b) through (d). Figure 3(b), which shows the microstructure of the sample annealed for 0.5 hours, reveals the formation of subgrains by nucleation (marked by -->), which occurs as a consequence of the re- covery process. In contrast, the annealing treatment of 1 hour produced a partially recrystallized structure, shown in Figure 3(c), that consisted of recovered subgrains and equiaxed recrystallized grains. Finally, the annealing time of 2 hours produced a fully recrystallized structure with equiaxed grains (Figure 3(d)). It has been suggested that a partially recrystallized microstructure suppresses hydrogen embrittlement,t20.211 and therefore, 1 hour was chosen as the recrystallization time in the present study.

The hardness of the alloy after the ordering treatment (380 VHN) is comparable to that after rolling (390 VHN). Heat treating the alloy at 500 ~ after thermomechanical processing would reduce the vacancy concentration and result in lower hardness. Moreover, the microstructures evolved after the ordering treatment (to be discussed in sec- tion III.C.3.) clearly indicate grain growth, the process that will reduce the dislocation density. If no other change takes place in the structure, these processes should lead to a de- crease in the hardness of the alloy. The formation of ordered phases, on the other hand, should result in increased hardness. In the present study, the fall in hardness due to the softening mechanisms appears to have been counter- acted by the formation of ordered phase(s). Therefore, the


1000

800

no 600

L +- ' 400

200

Rolted 1000~ 800~ 5 0 0 %

!

I I

i I

'// /

01/ I i I I I 0 0 . 0 5 0 .10

Strain (a)

(b)

(c)

(a)

Fig. 4---(a) Stress-strain behavior of the alloy rolled at 1000 ~ 800 ~ and 500 ~ Optical micrographs of the alloy rolled at (b) 1000 ~ (c) 800 ~ and (d) 500 ~

presence of ordered D O 3 phase after the ordering treatment is indirectly indicated, since the hardness after ordering is comparable to that after rolling.

C. Room-Temperature Tensile Behavior

1. As-rolled alloy Results of the room-temperature tensile tests conducted

after various TMT and post-TMT heat treatments of the Fe- 25A1 alloy are presented in Table II. The engineering stress- strain curves of samples rolled at 1000 ~ 800 ~ and 500

~ along with their microstructures are shown in Figure 4. It can be seen from Figure 4(a) that the ductility of the as- rolled alloy, as given by the plastic strain from the engineering stress-strain curve, was practically nil in all three cases and the fracture strength of the alloy gradually decreased with decreasing rolling temperature. This variation in fracture strength can be explained on the basis of microstructures of the alloy after its different rolling treatments.

Microstructures of the alloy rolled at 1000 ~ 800 ~

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 27A, OCTOBER 1996~2989

(a)

(b)

Fig. 5--Fracture morphology of the Fe-25Al alloy in (a) unetched condition and (b) etched condition.

and 500 ~ have been presented in Figures 4(b) through (d), respectively. As shown in Figure 4(b), the alloy rolled at 1000 ~ possessed numerous fine grains (marked as ---0 and regions of deformed elongated grains (marked as "l'), indicating the occurrence of some amount of dynamic re- covery (DRY) and dynamic recrystallization (DRX) processes when rolled to 80-pct thickness reduction at 1000 ~ As shown in Figure 4(c), after rolling at 800 ~ the alloy consisted of mainly coarse elongated grains along with very few recrystallized grains (marked as ,1,). In contrast to these structures that indicate the presence of dynamically occurring changes over clearly noticeable regions, the structure of the alloy rolled at 500 ~ as shown in Figure 4(d), consisted of highly deformed coarse grains with long stringers (marked as ,1,) and an extremely low volume frac- tion of fine subgrains (marked as ~ ) . Thus, it is clear that the grain size of the alloy varied considerably with the rolling temperature. The variation in the strength of the as- rolled alloy can thus be understood in terms of the Hall- Petch relationship.

The fracture in all three samples occurred in a transgran- ular or brittle cleavagelike manner, as seen in the SEM fractograph of Figure 5(a). The fractograph shows a river line pattern of curvilinear shape (marked as 1") running par- allel to each other and possibly formed as a consequence of convergence of irregular cleavage faces, t221 Such river line patterns are often formed as a result of the presence of screw dislocations, and it is interesting to note that Morris

and Leboeuf have found that subgrain boundaries of iron aluminides are generally composed of (100) dislocations that are of screw orientation.t231

Emerging points of dislocations can sometimes be observed by suitable etching.t241 Since the lattice around a dislocation line is in a highly strained state, it is preferentially attacked by the etchant. The fractograph obtained from the sample rolled at 1000 ~ after its etching is shown in Figure 5(b). This figure shows that most of the etch pits are of rectangular shape, indicating that the cleavage planes are of cube type, e.g., (100).

Observation of (100)-type cleavage planes under conditions of embrittlement can be explained in the following manner. Mobile dislocations trap hydrogen and carry it to potential flaw sites such as crack tips, inclusions, and voids. In intermetallic structures, the dislocations are generally in the form of partials. The Fe3AI has a high stacking fault energy (SFE) value, t25j which is of the order of 500 ergs/cm 2. Thus, it forms partials that have a low splitting width. These hydrogen-containing partials glide on slip planes in the ordered B2 and DO3 structures of iron aluminides and can be locked up as immobile (100) dislocations by the Freidel mechanism.t2=1 Moreover, hydrogen promotes the formation of locked dislocations, and twice the amount of normal hydrogen can be absorbed by locked (100) dislocations, t261 According to most of the accepted theories of hydrogen embrittlement, hydrogen should ex- ceed a critical concentration at potential flaw sites to cause embrittlement. Critical hydrogen concentration is a function of stresses generated by hydrogen occupying tetrahedral interstitial sites on the (100) planes of the ordered B2 and D O 3 structures. The actual number of tetrahedral interstitial sites would depend on the crystal structure (B2 or DO~), but there are enough interstitial sites, which leads to hydrogen concentration and finally crack initiation by the decohesion mechanism, r~~

2. Rolled and recrystallized alloy It has already been stated that for the recrystallization

anneal, a temperature of 750 ~ was chosen and in order to obtain a partially recrystallized structure an annealing time of 1 hour was selected. Figure 6(a) shows engineering stress-strain curves for the alloy rolled at three different temperatures and then recrystallized at 750 ~ for 1 hour. A significant change in the fracture strength was observed for the alloy rolled at different temperatures. While the alloy rolled at 500 ~ exhibited the highest fracture strength of 730 MPa, those rolled at 800 ~ and 1000 ~ showed lower fracture strengths of 570 and 460 MPa, respectively.

The variation of fracture strength of the recrystallized alloys after rolling at different temperatures can be explained in terms of their microstructures, which are presented in Figures 6(b) through (d). As shown in Figure 6(b), the alloy rolled at 1000 ~ possessed a coarse grain structure after its recrystallization, because in the as-rolled condition, it had undergone dynamic recrystallization and the subgrains formed during this process underwent grain growth during recrystallization annealing. On the other hand, the specimen rolled at 800 ~ exhibited a mixture of

large subgrains (marked as ~) and equiaxed nucleated grains

(marked as q'). However, these equiaxed grains have not undergone grain growth, as in the specimen rolled at 1000 ~


800

600

400

200

RecrystaUized 1000~ / 800~ / 500"C /

#

- / /

./ / / /

/ / /

, I / / / Z'"

f / /

0 I I i 0 0.02 0.04 0.06 0.08

Strain (a)

(b)

(c)

(a0

Fig. 6--(a) Stress-strain behavior of recrystallized alloy initially rolled at 1000 ~ 800 ~ and 500 ~ Optical micrographs of the alloy (b) rolled at 1000 ~ and recrystallized, (c) rolled at 800 ~ and recrystallized, and (d) rolled at 500 ~ and recrystallized.

Finally, since dynamically occurring changes were almost insignificant in the case of the specimen rolled at 500 ~ small subgrains were nucleated during annealing of the specimen and this resulted in a fine grain structure. The large amount of stored strain energy arising because of the thickness reduction of 80 pct imparted to the specimen during rolling at 500 ~ acted as the driving force for the formation of a large number of nuclei, resulting in a fine grain structure. Thus, the alloy rolled at 500 ~ and then recrystallized exhibited the highest fracture strength as a result of its fine grain structure.

Figure 6(a) also shows that the highest ductility was obtained in the case of the alloy rolled at 1000 ~ and then

recrystallized. This could possibly be attributed to the low grain-boundary area as revealed by its microstructure in this condition. McKamey and Pierce have suggested earlier that the reduction in grain-boundary area leads to an increase in ductility, as it suppresses hydrogen diffusion in the lattice.t201

3. Rolled and ordered alloy The mechanical properties of the rolled Fe-25A1 alloy

after an ordering treatment of 100 hours at 500 ~ are presented in Table II and the engineering stress-strain curves are shown in Figure 7(a). It is clear from these results that the ductility was practically zero in all three cases. Frac- tographs of the samples showed a brittle cleavage type of


600

~. 400

200

Ordered looo~ /

_ _ eo0~ / _- .- 500*////

/,/ // /" / /

- / / /

/i//1 , y" I"

0 I I 0 0.04 0.08 0.12

Strain (a)

(b)

(c)

(d) Fig. 7--(a) Stress-strain behavior of the alloy ordered at 500 ~ for 100 h after initially being rolled at 1000 ~ 800 ~ and 500 ~ Optical micrographs of the alloy (b) rolled at 1000 ~ and ordered, (c) rolled at 800 ~ and ordered, and (d) rolled at 500 ~ and ordered.

fracture, similar to that found in the as-rolled and rolled and recrystallized alloys.

Microstructures of specimen rolled at three different temperatures and subsequently ordered are shown in Figures 7(b) through (d). The alloy rolled at 500 ~ had the highest stored energy after its deformation processing and possessed the greatest driving force for the ordering treatment. Therefore, it underwent grain growth to a greater extent. The lower fracture strength of the specimen rolled at 500 ~ could be attributed to its coarse grain structure. As the stored strain energy was less in the case of specimens rolled

at 800 ~ and 1000 ~ the extent of grain growth was found to be reduced in these specimens. In fact, in the case of the specimen rolled at 1000 ~ some elongated grains persisted even after the ordering treatment. The mechanical properties obtained after ordering could thus be explained based on the microstructure.

IV. CONCLUSIONS

1. The Fe-25A1 alloy exhibited the highest hardness in the as-rolled condition compared to recrystallized and or-


dered conditions. This has been explained by cooling rate effects.

2. The presence of D O 3 phase in Fe-25A1 alloy after ordering treatment has been confirmed by (440) peak broadening in the XRD pattern.

3. The alloy rolled at 1000 ~ showed greater fracture strength in the rolled condition compared to the alloy rolled at 800 ~ and 500 ~ because of its finer grain size. This has been explained by considering dynamic recrystallization occurring at these temperatures and the resulting microstructures.

4. The alloy rolled at 500 ~ and then recrystallized at 750 ~ for 1 hour showed the highest fracture strength compared to those rolled at 800 ~ and 1000 ~ and subsequently recrystallized. However, the alloy rolled at 1000 ~ and then recrystallized at 750 ~ for 1 hour exhibited the highest ductility. This behavior has been ascribed to the degree of recrystallization and the microstructure. Partially recrystallized microstructures pre- vent hydrogen ingress through grain boundaries and thus minimize hydrogen embrittlement.

5. The alloy rolled at 1000 ~ and then ordered at 500 ~ for 100 hours showed the highest fracture strength compared to those rolled at 800 ~ and 500 ~ and subsequently ordered. This behavior is attributable to its smaller grain size. The variation in grain size after ordering has been discussed.

6. Fracture morphologies were found to be of brittle fracture. The fractographs revealed river flow patterns char- acteristic of cleavage type of fracture. Etching study indicated that the cleavage facets of the fractured surface were of cube orientation, e.g., of (100) type. This has been explained based on the decohesion mechanism of hydrogen embrittlement.

ACKNOWLEDGMENTS

The authors gratefully acknowledge Dr. D. Banerjee and Dr. A.K. Gogia of Defence Metallurgical Research Labo- ratory, India for providing the alloy used in the study.

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14, pp. 8-13. 15. H.J. McQueen and G.C. Kuczynski: Trans. TMS-AIME, 1959, vol.

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21. V.K. Sikka, S. Vishwanathan, and C.G. McKamey: in Structural Intermetallics, R. Daroha, J.J. Kewandoski, C.T. Liu, P.L. Martin, and M.B. Nathal, eds., TMS, Warrendale, PA, 1993, pp. 483-91.

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