Influence of processing parameters on dislocation substractures in deformed-and-recovered FeVNi solid solutions

Materials Science and Engineering, 63 (1984) 239-250

239

Influence of P r o c e s s i n g Parameters on Dislocation Substructures in Deformedand-recovered F e - V - N i Solid Solutions E. R. BRADLEY* and D. H. POLONIS

Department of Materials Science and Engineering, University of Washington, Seattle, WA 98195 (U.S.A.) (Received August 22, 1983i

SUMMARY

T w o modes o f deformation have been employed to modify the dislocation substructures in ferritic F e - V - N i alloys. The alloys were deformed to 30% and 75% reduction in area and annealed at temperatures from 300 to 500 °C to produce a variety o f subgrain sizes, shapes and misorientations. Deformation by swaging produces a uniform distribution o f nearly equiaxed subgrains, whereas rolling results in deformation bands, leading to non-uniform dislocation substructures. The subgrain size in the swaged-and-recovered alloys increases with increasing recovery temperature; increasing the nickel content reduces the limiting subgrain size in the deformed-and-recovered condition. The average subgrain misorientation increases with increasing amounts o f deformation, ranging from 1.05 ° after 30% reduction to 2.5 ° for specimens deformed at 75%. Increasing the nickel content from 1% to 3% results in increased thermal stability o f the subgrains, thereby p r o m o t i n g smaller subgrain sizes in the higher nickel composition alloys. This observation is consistent with the enhancement o f cross-slip and the restriction o f dynamic recovery, both o f which are expected with increasing nickel con ten t.

1. INTRODUCTION The grain refinement of ferritic steels is of great interest because it leads to an increase in the yield strength and a concurrent decrease in the ductile-to-brittle transition temperature [1-3]. However, there are practical limitations to the degree of grain refinement attainable *Present address: Battelle Pacific N o r t h w e s t

tories, Richland, WA, U.S.A. 0025-5416/84/$3.00

Labora-

by normal thermal or thermomechanical processing methods. The introduction of dislocation subboundaries into a fine-grained material provides an alternative means of extending the beneficial effect of grain boundaries, but the relationships between subgrain structure and the strength and fracture characteristics of iron need to be defined more clearly before this strengthening method can be used to its full potential. The subgrain size reduction observed during tensile deformation is expected to differ in materials deformed by more complex modes such as rolling, swaging or drawing. In the latter cases the subgrains continue to decrease in size with increasing deformation as a result of the constraints imposed by the deformation mode; also, the subgrains are distorted in a response to the overall shape change of the specimen [4, 5]. Nickel additions to ~-Fe have been reported to increase cross-slip and to improve the strength and low temperature fracture resistance [6, 7]. Consequently, small additions of nickel to ferritic alloys can be expected to m o d i f y the dislocation substructures produced by deformation, but this effect has not been investigated previously. Small amounts of vanadium remove interstitial carbon and nitrogen from solution, thereby improving the fracture resistance at low temperatures. In the present paper the dislocation substructures formed in ferritic solid solutions of Fe-V-Ni as a result of deformation by two differing modes, namely swaging and rolling, followed by recovery annealing treatments are examined. 2. DISLOCATION SUBSTRUCTURES IN IRON The relationship between dislocation density and plastic strain in iron is approximately © Elsevier Sequoia/Printed in The Netherlands

240 linear during the initial stages of deformation, b u t the rate of increase is reduced at higher strains [8, 9]. The probability of dislocation annihilation by dynamic recovery increases with increasing dislocation density, thereby accounting for the departure from linearity at high strains [10]. The average density of dislocations at a given strain is inversely related to the initial grain size [8, 9], because of an increase in the number of potential dislocation sources in fine-grained materials [ 11]. For strains less than 1%, iron specimens deformed in tension at room temperature exhibit arrays of relatively straight dislocations [9]. With increasing deformation, the dislocations become jogged and interactions between dislocations lead to the formation of dislocation clusters. These clusters of high dislocation density continue to develop with increasing strain and eventually join to form a subgrain structure. The subgrain boundaries consist of tangled networks of dislocations surrounding volumes having relatively low densities of dislocations. The subgrain size decreases initially with strain and then remains essentially constant or decreases at a reduced rate, depending on the mode of deformation [ 5, 9]. The misorientation between subgrains increases with deformation because of the accumulation of dislocations in the subgrain boundaries. The processes involved in the recovery of dislocation substructures produced by deformation are very temperature dependent. A t low temperatures (T ~ 0.2Tin where Tm is the melting temperature) the removal and redistribution of point defects occur with little change in the dislocation substructure [8, 12]. As the recovery temperature is increased, the subgrain size increases while the overall dislocation density is decreased [8, 9, 13]. At temperatures approaching 0.4Tin, large-scale rearrangements produce low energy dislocation configurations in the subgrain boundaries [9, 13, 14]. The development of a uniform and stable dislocation substructure is an important objective in utilizing mechanical-thermal processing for the development of high strength materials. Two modes of deformation, rolling and swaging, were used to modify the dislocation subgrain structure at either 30% or 75% reduction in area followed by annealing at temperatures in the range 3 0 0 - 5 0 0 °C.

3. EXPERIMENTAL PROCEDURES Ternary F e - V - N i alloys were prepared by arc melting in purified argon using a Centorr single-electrode arc-melting furnace. High purity {99.9%) nickel and vanadium were added to Ferrovac E grade iron. The arcmelted buttons were inverted and remelted five times to ensure homogeneity. The circular buttons were cut to appropriate sizes, remelted and cast into either rod or rectangular shapes for swaging or rolling. The cast forms were annealed at 1000 °C for 3 h in either high vacuum or purified helium followed by furnace cooling. This heat treatment served as a final homogenization process that transformed the cast structure and produced an equiaxed grain morphology. The cast-and-homogenized alloys were initially swaged or rolled to the appropriate size for subsequent processing and then heat treated at 1000 °C for 30 min, followed by furnace cooling. The resulting grain sizes were 40 pm and 25 pm respectively for the Fe1%V.-1%Ni and F e - l % V - 3 % N i alloys. Final deformation consisted of swaging to 30% or 75% reduction in area or rolling to 30% or 75% reduction in thickness, thereby producing approximately equivalent amounts of strain by both modes of deformation. Subsequent annealing treatments were carried out as summarized in Table 1. Transmission electron microscopy was performed using a JEM 200 A microscope with a large-angle goniometer stage and an attachment for high resolution dark field microscopy. The latter attachment was especially useful for centering the transmitted electron beam which was severely displaced because of the ferromagnetic characteristics of the alloys. Thin foils for transmission electron microscopy were prepared from transverse sections, approximately 0.5 mm thick, that were cut from the swaged rods 3 mm in diameter and the rolled sheet 1 mm thick. These sections were mechanically thinned to 0.1 mm by hand grinding and then they were electrolytically thinned to electron transparency by a two-step procedure, i.e. both surfaces were dimpled first and then electropolished to perforation. The disc-shaped specimens from the swaged rods were dimpled in a twin-jet electropolishing unit at an applied voltage of 35 V using 10% perchloric acid in butyl

241 TABLE 1 Summary of the recovery annealing treatments Annealing Temperature

Time

(°C)

(h)

325 370 400 440 500 540

Rolled alloys (all)

S w a g e d alloys

6 6 6 6 2 1

Fe-1%V-1%Ni

F e - 1 % V - 3 %Ni

30% r e d u c t i o n in area

75% r e d u c t i o n in area

30% r e d u c t i o n in area

75% r e d u c t i o n in area

× X X ×

X X × × × ×

X X × × ×

X

alcohol as the electrolyte. An apparatus similar to that described by DuBose and Jones [15] was used for the rolled alloys because the specimens were too small for the twin-jet unit. A 20~ perchloric acid in butyl alcohol electrolyte was used and the applied voltage was 150 V for these specimens. The dimpled specimens were then electropolished at a potential of 10 V in a 10% perchloric acid in butyl alcohol solution and cooled to 0 °C by an ice-water bath. The technique described by DuBose and Jones [15] was used for detecting the initial perforation. The subgrain sizes were measured from photographic enlargements by an intercept technique using a circular test line 75 mm in diameter inscribed on a transparent sheet. The test line was randomly positioned on the electron micrographs and the subgrain boundaries intersecting the test line were counted. At least five electron micrographs representing different regions of the samples were used with at least 20 individual measurements being made on each micrograph. The average intercept length was determined by the relation 1 t -- _ NL

(1)

where NL is the average number of intersections per unit length of test line. The subgrain shape was characterized by the Smith and G u t t m a n [16] shape parameter: 7r shape parameter -- N t 2

×

X × × X

×

where N is the number of subgrains per unit area and t is the average intercept length. The misorientation across subgrain boundaries was determined from electron diffraction patterns of adjacent subgrains by measuring the shift in the Kikuchi line pattern. To utilize the technique, it was necessary to use a relatively thick (about 2000 A) region of the foil and a small selective diffraction aperture so that the diffracted intensity was entirely from the individual subgrain being examined. This latter requirement limited the minimum subgrain diameter that could be examined, i.e. approximately 0.7 gm for the electron microscope used in the present work. The shift in the Kikuchi lines was determined by overlaying the electron diffraction negatives from adjacent subgrains such that the two Kikuchi patterns were in coincidence and then measuring the displacement of the diffracted spot pattern. The misorientation between the two subgrains was then calculated from the relation [17] X 0 -- - L

where 0 is t h e angular m i s o r i e n t a t i o n in radians, X is t h e d i s p l a c e m e n t o f t h e Kikuchi lines and L is t h e effective c a m e r a l e n g t h o f t h e camera.

4. EXPERIMENTAL RESULTS

(2)

(3)

The substructures produced by the combined mechanical-thermal treatments were examined by transmission electron microscopy. Comparisons were made between the

242 subgrain p a r a m e t e r s in c o l d - w o r k e d F e - l % V l % N i and F e - l % V - 3 % N i alloy specimens t h a t had been annealed at t e m p e r a t u r e s b e t w e e n 300 and 550 °C. The m e t h o d s used t o determine the size, shape and m i s o r i e n t a t i o n o f the subgrains were as described previously.

4.1. Substructure in swaged alloys Well-defined subgrain structures were f o r m e d in all the swaged-and-recovered alloys, as illustrated in Fig. 1; this m i c r o s t r u c t u r e o c c u r r e d i n d e p e n d e n t l y of the a m o u n t o f d e f o r m a t i o n , the alloy c o m p o s i t i o n and the r e c o v e r y t e m p e r a t u r e . The structures were u n i f o r m t h r o u g h o u t the transverse sections e x c e p t for a few small areas in which the subgrains were larger t h a n average and the disl o c a t i o n d e n s i t y in t h e b o u n d a r i e s was lower, as s h o w n in Fig. 2. Such areas were observed t o some e x t e n t in all t h e swaged-and-recovered specimens b u t were m o r e f r e q u e n t in specimens swaged to 30% r e d u c t i o n in area. The dislocations were c o n f i n e d primarily t o subgrain b o u n d a r i e s with a fairly low disloca-

tion d e n s i t y within subgrains. F o r specimens annealed at 400 °C or lower, the b o u n d a r i e s consisted o f tangled arrays o f dislocations, as s h o w n in Fig. 2. Increasing the r e c o v e r y annealing t e m p e r a t u r e t o above 400 °C resulted in extensive r e a r r a n g e m e n t o f the b o u n d a r y dislocations and p o l y g o n i z e d b o u n d a r i e s were o f t e n resolvable. Most o f the p o l y g o n i z e d b o u n d a r i e s consisted o f cross-meshed or h e x a g o n a l dislocation n e t w o r k s , as s h o w n in Fig. 3, suggesting t h a t d e f o r m a t i o n b y swaging p r o d u c e s subgrain b o u n d a r i e s with p r e d o m i n a n t l y twist b o u n d a r y characteristics. I t was n o t possible to resolve dislocation n e t w o r k s in b o u n d a r i e s having m i s o r i e n t a t i o n s greater t h a n a b o u t 1 °, so their e x a c t n a t u r e was n o t identified. Table 2 is a c o m p i l a t i o n o f the transverse subgrain p a r a m e t e r s m e a s u r e d f r o m transmission e l e c t r o n m i c r o g r a p h s of the swagedand-recovered alloys. D a t a are given for t w o groups o f F e - 1 % V - l % N i alloy specimens t h a t were swaged to 75% r e d u c t i o n in area u n d e r d i f f e r e n t settings o f the swaging dies. The

TABLE 2 Subgrain parameters of the swaged-and-recovered Fe-l%V~1%Ni and Fe-1%V-3%Ni alloys

Alloy

Fe-l%V-l%Ni

Fe-l%V-3%Ni

aGroup A specimens. bGroup B specimens.

Cold working

Recovery temperature

Subgrain size

Subgrain misorientation

(%)

(°C)

(pm)

(deg)

Shape parameter Tr/Nt2

30 20 30 30

370 400 440 500

0.67 0.67 0.73 0.81

1.04 -1.04 1.08

5.1 4.5 4.8 4.2

75 a 75 a 75 a 75 a 75 a

325 370 415 440 540

0.58 0.61 0.73 0.83 1.01

---2.10 2.08

4.4 4.4 4.2 4.2 4.5

75 b 75 b 755 755

370 400 440 500

0.57 0.64 0.68 0.87

---2.50

4.2 3.8 4.0 3.7

30 30 30 30 3O

370 400 440 500 540

0.40 0.46 0.51 0.53 0.57

----1.54

4.8 5.3 3.3 4.7 4.5

75 75 75

370 440 540

0.40 0.50 0.62

----

4.4 4.7 4.2

243

Fig. 1. Transmission electron micrograph showing the uniform subgrain structure from a transverse section of an F e - l % V - 1 % N i alloy (group B) specimen swaged to 75% reduction in area and annealed at 500 °C.

Fig. 2. Transmission electron micrograph showing enlarged subgrains in an isolated region of an F e - l % V - l % N i alloy specimen swaged to 30% reduction in area and annealed at 370 °C.

244

j

0.8

Of

03 0.6

OI Q ~

0.5 --

E

I

0.4

(a)

I

I

o"

_z l.O -

0.9

0.8 0.7 0.6

Fig. 3. Transmission electron micrograph showing cross-meshed and hexagonal dislocation structures in the subboundaries of an F e - l % V - 1 % N i alloy swaged to 30% reduction in area and annealed at 440 °C.

(b) 300

m~

I

i

400

500

600

RECOVERYTEMPERATURE(°C)

group A specimens were swaged under conditions where the displacements of the dies in the open position were less than for the remaining specimens. At least five electron micrographs and 100 individual measurements were used to determine the average intercept subgrain sizes given in Table 2. The data followed normal distributions and, although there was some overlap in the size distributions for adjacent temperature levels, the average subgrain size was clearly defined in each case. The average size of the subgrains increased with increasing recovery temperature for a given alloy and deformation condition, as shown in Fig. 4. The rate of increase in subgrain size with increasing recovery temperature depended on both the a m o u n t of deformation and the alloy composition, as indicated by the slopes of the curves in Fig. 4. A linear: regression analysis revealed an increase by a factor of 2 in the slopes for the F e - I % V - I % N i specimens when the a m o u n t of prior deformation was increased from 30% to 75% reduction in area. A smaller increase, about 20%, is observed for the Fe-1%V-3%Ni data; also the Fe-l%V-3%Ni alloys exhibit lower slopes than the Fe-1%V-l%Ni alloys deformed equally. The subgrain size was smaller in the Fe-l%V-3%Ni alloy specimens at all recovery temperatures. The misorientations given in Ta-ble 2 were determined by measuring the translation of

Fig. 4. Subgrain size as a function of the recovery annealing temperature in the swaged-and-recovered F e - V - N i alloys: (a) 30% reduction in area (©, FeI % V - I % N i ; e, Fe- l %V - 3 %N i ) ; (b) 75% reduction in area (~, Fe-l%V-1%Ni, group A; A, F e- 1 %V l%Ni, group B ; . , Fe-l%V-3%Ni).

the Kikuchi patterns obtained from adjacent subgrains. The measurements of the misorientation by this method are limited to subgrains having diameters greater than about 0.7/~m and with misorientations less than about 4 °. These limitations permitted the misorientation to be determined as a function of the recovery annealing temperature only in the Fe-l%V-1%Ni alloy specimens swaged to 30% reduction in area. The average misorientation was found to be 1.04 ° in the specimen annealed at 370 °C and 1.08 ° for the specimen annealed at 500 °C. These values are based on approximately 20 individual measurements on each specimen. The small difference in misorientation is not statistically significant. As shown in Table 2, increasing the deformation from 30% to 75% reduction in area increases the average misorientation across the subgrain boundaries to greater than 2 ° in the F e - I % V - I % N i alloys. Attempts to measure misorientations across subgrain boundaries in the Fe-l%V-3%Ni alloy specimens were mostly unsuccessful because of the smaller subgrain size in this alloy. Six successful measurements were made on the specimen swaged to 30% reduction in

245 area and annealed at 540 °C. The average of these measurements was 1.54 °, a value significantly larger than that for the F e - I % V - I % N i alloys swaged to 30% reduction in area. The shapes of the subgrains in transverse sections were characterized by utilizing the shape parameter of Smith and G u t t m a n [16] : 7~ shape parameter = Nt 2 where N is the number of subgrains per unit area and t is the mean intercept length between subgrain walls. For reference, the shape parameter is 4 for circles, 4.4 for regular hexagons, 5.1 for squares and 5.7 for 2 X 1 rectangles. From Table 2 it can be seen that the shape parameters for the swaged-andrecovered alloys are all in the range 3.7-5.3, indicating the equiaxed nature of the subgrain structure. The shape parameters also appear to be rather insensitive to the recovery temperature, to the a m o u n t of deformation or to the alloy content. Figure 5 shows the longitudinal subgrain structure from the Fe-1%V-l%Ni (group B) alloy deformed to 75% reduction in area and annealed at 500 °C. The subgrains were aligned along the axis of the rod with varying degrees of elongation. The shape parameter was 5.9, which suggests 2 X 1 rectangles. Mea-

surements on individual subgrains are consistent with this subgrain shape.

4.2. Substructures in rolled alloys In contrast with substructures in the swaged alloys, uniform subgrain structures were not observed in the rolled-and-recovered alloys. The dislocation substructures consisted of some regions with elongated subgrains and other regions with a high dislocation density but having no apparent subgrain formation. Examples of the two microstructure types are shown in the electron micrographs of Fig. 6, representing a transverse section of Fe-l%V-3%Ni alloys rolled to 75% reduction in area and annealed at 440 °C. Similar microstructures were observed in all the rolled-and-recovered specimens, indicating that the deformation mode was responsible for the non-uniform substructures. The wide variations in substructure from one region to another allowed only qualitative comparisons of dislocation substructures in the rolled-and-recovered alloys. In general, the microstructures varied in accordance with the expected changes resulting from the a m o u n t of deformation and the annealing temperatures employed. Increasing the a m o u n t of deformation produced more regions with

Fig. 5. Transmission electron micrograph showing the subgrain structure from a longitudinal section of an F e 1%V-1%Ni alloy specimen swaged to 75% reduction in area and annealed at 500 °C.

246 the Fe-1%V-3%Ni alloys they were observed only after annealing at 500 °C. Figure 7 shows a region in the F e - I % V - I % N i alloy annealed at 500 °C where recrystallization had initiated in two parallel bands separated by a highly distorted subgrain structure. Specimens rolled to 30% reduction in area showed no evidence of recrystallization after annealing at 500 °C, which suggests that a critical amount of deformation is needed to initiate recrystallization at these low temperatures. 5.: DISCUSSION OF RESULTS

• I,fl

! ~iii~'

"

"Fig. 6. Transmission electron micrographs showing the diverse microstructures in an Fe-l%V-3%Ni alloy rolled to 75% reduction in area and annealed at 440 °C: (a) elongated subgrains; (b) high dislocation density with no apparent subgrain structure.

elongated subgrains and the subgrains became more clearly defined because of higher boundary misorientations. A general decrease in dislocation density within the elongated subgrains was observed with increasing recovery temperature. The width of the elongated subgrains varied between 0.2 and 0.5 pm in allthe samples and the wide variations in substructure made quantitative comparisons of the subgrain size meaningless. In addition to the diverse dislocation subgrain structures, specimens rolled to 75% reduction in area and annealed at the higher recovery temperatures showed isolated recrystallized regions corresponding to less than 1% of the area examined. These regions existed in the F e - I % V - I % N i alloy specimens annealed at both 440 and 500 °C, whereas in

5.1. Substructures in the swaged-andrecovered alloys The shape of the dislocation subgrains in the swaged alloys is in general agreement with that expected for this type of deformation process. The axially symmetric nature of the deformation is expected to produce equiaxed subgrains in the transverse sections and elongated subgrains in the longitudinal direction. This microstructure is appropriate for quantitative characterization of the subgrain size and misorientation, thereby providing a basis for evaluating the effects of mechanicalthermal processing variables on these important strengthening parameters. Previous studies have shown that the subgrain size initially decreases rapidly with increasing strain and then either saturates at a lower limit [9] or continues to decrease, but at a much lower rate [5]. These changes in the observed subgrain size with deformation are due to a direct competition between the processes that reduce the subgrain size, such as subgrain division by dislocation interactions, or the overall shape change of the specimen and dynamic recovery, whereby subboundary migration and subgrain coalescence tend to increase the average subgrain size. Hence, the processes leading to a decrease in subgrain size predominate during the initial stages of deformation with dynamic recovery becoming increasingly more important as deformation proceeds. The role of dynamic recovery in the swaged alloys can be assessed by comparing the measured subgrain sizes at the two levels of deformation with the sizes predicted for axially symmetric deformation. For axially symmetric flow the reduction in the mean transverse linear intercept subgrain size with strain

247

Fig. 7. Transmission electron micrograph showing the initial stages of recrystallization of an Fe-1%V-1%Ni alloy rolled to 75% reduction in area and annealed at 500 °C. can be expressed as t=tiexp

--

(4)

where t and ti are the mean intercept sizes at true strain values of e and ei respectively. The true strain for the swaged rods can be estimated from the relation

where Ao and At are the initial and final crosssectional areas of the rods. Using these relations, a decrease in the mean linear intercept subgrain size o f a p p r o x i m a t e l y 40% is expected when the d e f o r m a t i o n is increased from 30% to 75% r e d u c t i o n in area. The measured subgrain size for specimens d e f o r m e d to 30% r e duc t i on in area and annealed at 370 °C was used as a basis for applying eqn. (4). The subgrain sizes predicted by

the equation for specimens d e f o r m e d t o 75% reduction in area are 0.40 pm and 0.24 pm for the F e - I % V - I % N i and the F e - l % V - 3 % N i alloys respectively. The m i ni m um subgrain sizes measured in these alloys are 0.57 pm and 0.40 pm respectively; the differences, of approxi m at el y 67% between the predicted and measured sizes, can be accounted for if dynamic recovery occurs during the developm e n t of the dislocation substructures. It may be argued that this difference can result from a higher rate of subgrain growth in the m ore heavily d e f o r m e d materials. However, when the experimental data are linearly extrapolated to 300 °C where subgrain growth is n o t expected [18], the extrapolated subgrain sizes remain greater than the predicted sizes by approxi m at el y 15% for the F e - 1 % V l%Ni alloy and by 40% for the F e - l % V - 3 % N i alloy. These significant differences between the predicted and extrapolated experimental values lend credence to the argument that

248 dynamic recovery is influential in limiting the subgraln size. The concept of a limiting subgrain size following combined mechanical-thermal processing is especially evident in the specimens recovered at the higher temperatures where the subgrain size is larger in the more heavily deformed alloys (see Table 2). Under these conditions the larger growth rate in the heavily deformed alloys exceeds the tendency for a decrease in size due to increasing deformation. Thus, the competition between growth rate and subgrain refinement due to increased plastic strain results in a minimum subgrain size that can be achieved by combined mechanical-thermal processing. The data indicate a minimum subgrain size of about 0.55 #m for the F e - 1 % V - l % N i alloy and 0.40 pm for the Fe-1%V-3%Ni alloy. The smaller subgrain size in the F e - l % V 3%Ni alloy can be attributed to several factors, including (1) differences in the initial grain size, {2) differences in cell formation and subdivision and (3) differences in dynamic recovery. The initial grain diameter was a factor of almost 2 smaller in the F e - l % V - 3 % N i alloy than in the F e - l % V - 1 % N i alloy; this is expected to result in an increased dislocation density during the initial stages of deformation [8, 9]. Holt's [19] theory of subgrain formation suggests that the subgrain size developed during deformation is inversely proportional to the dislocation density. Consequently, the smaller subgrain size developed during the initial stages of deformation in the F e - l % V - 3 % N i alloys is probably due to the increased dislocation density associated with the smaller grain size. The specific effects of nickel additions on the production and interaction of dislocations in dilute ferritic alloys are not known and these are important considerations in assessing both the effects of nickel on the formation of a subgrain structure and the subsequent subdivisions of the subgrains with additional deformation. Subgrain formation is related to the ability of dislocations to cross slip [ 20]; since the addition of nickel to iron enhances cross-slip in dilute ferritic alloys [6, 7], it is reasonable to expect that subgrain formation and subdivision will be enhanced b y nickel additions. This would also result in a decrease in the limiting subgrain size in the Fe-1%V-

3%Ni alloy compared with that in the alloy containing only 1% Ni. Nickel additions could also reduce the limiting subgrain size by restricting dynamic recovery. Solute additions are known to reduce the thermal mobility of boundaries [ 21 ] as well as the movement of the individual dislocations, and a similar effect is expected for the athermal migration of cell boundaries during deformation. This would decrease the rate of subgrain coalescence during dynamic recovery and thereby reduce the limiting subgrain size at large deformations. The effects of the initial grain size and cross-slip processes on the subgrain size are expected to be most important during the initial stages of deformation. Conversely, the reduced boundary mobility should predominate at a later stage of deformation when dynamic recovery becomes an important process. Dynamic recovery in the swaged alloys has already been established as an important factor and therefore the smaller subgrain size in the F e - l % V - 3 % N i alloy is most probably due to the restricted mobility of the subgrain boundaries. The restricted subboundary mobility in the F e - l % V - 3 % N i alloy is also shown in Fig. 5 where the temperature dependence of the subgrain size is less than that of the F e - l % V - l % N i alloys subjected to comparable amounts of deformation. In contrast with the subgrain size the misorientation across the subgrain boundaries increased by about a factor of 2 with an increase in deformation from 30% to 75% reduction in area (see Table 2). Langford and Cohen [22] have reported a linear increase in the average misorientation (up to 10 °) with strain in ~-Fe deformed by wire drawing to large reductions in area (greater than 99%). The results from the swaged alloys are in good agreement with their measurements at comparable strains and it is reasonable to expect that the subgrain misorientations will continue to increase with increasing amounts of swaging deformation. A comparison of the parameters in Table 2 shows that the group A F e - l % V - 1 % N i specimens have a larger subgrain size for equivalent recovery temperatures and a smaller misorientation than the group B specimens which were deformed equivalent amounts. These differences can be accounted for, at least in part, by the nature of the flow processes

249 accompanying swaging in both cases. The group A specimens were swaged under conditions where the displacement of the two dies in the open position was less than for the remaining specimens. Swaging die cavities are slightly oval to prevent sticking; the reduced clearance between the dies when the group A specimens were swaged produced a torque on the material during the rotation of the dies. The specimen rotation leads to a spiral deformation pattern and the reduced clearance slows the feed into the dies, thereby reducing the rate of deformation. Michalak and Cuddy [23] have observed a decrease in cell size with increasing deformation rate, which is in accordance with the difference between the group A and group B specimens. However, other factors, such as the redundant and frictional work expended during the deformation process, are also expected to influence the microstructures. 5.2. Substructures in the rolled-and-recovered alloys The ribbon-shaped subgrains observed in the rolled alloys (Figs. 6 and 7) reflect the overall shape change of the specimen. This subgrain shape is c o m m o n l y observed in coldrolled materials [18]. However, the transverse subgrain dimensions did not change appreciably with increasing deformation, which is consistent with the important role of dynamic recovery in the development of the substructures. The wide variations in microstructure from one region to another reflect the inhomogeneous nature of deformation produced by rolling. Similar microstructures to those shown in Fig. 6 have been reported by other investigators [24-26] to be associated with the deformation bands which form in both single-crystal and polycrystalline ferritic alloys. Single crystals of Si-Fe alloys exhibit a high density of uniformly distributed dislocations within deformation bands and also elongated subgrains in a transition zone separating adjacent deformation bands [25]. The micrographs shown in Fig. 6 are consistent with these features and, since similar microstructures were observed in all the rolled-and-recovered specimens examined, it is concluded that the formation of deformation bands is responsible for the inhomogeneous microstructures.

Roberts and Jolley [27] reported a uniform distribution of equiaxed subgrains in coldrolled (~-Fe after annealing in the same temperature range as that employed in the present investigation. The equiaxed subgrains developed from initially elongated subgrains during recovery annealing. A similar change in subgrain shape was not observed in the annealed alloys in the present work, a difference that could be attributed to alloy additions restricting the mobility of small-angle boundaries. However, at the higher annealing temperature employed in the present work the elongated substructures persisted in the ferritic alloys even though recrystallization was initiated at preferred sites. Such microstructural changes due to annealing are expected in materials exhibiting deformation bands [12, 25, 28], thereby suggesting that deformation bands may not have formed in the materials examined by Roberts and Jolley. The difference between the dislocation substructures observed in the two investigations can be explained in terms of different textures in the starting materials because the formation of deformation bands is known to depend on the orientation of crystals relative to the rolling plane and the rolling direction [28]. 6. CONCLUSIONS (1) The dislocation substructure produced b y mechanical-thermal processing is very dependent on the mode of deformation. A uniform distribution of nearly equiaxed subgrains is produced by swaging, whereas rolling produces deformation bands which result in non-uniform dislocation substructures. (2) The subgrain size in swaged-andrecovered alloys increases with increasing recovery temperature. The increase is more pronounced in the alloys deformed to 75% reduction in area and, when combined with dynamic recovery during the deformation, results in a minimum subgrain size that can be achieved b y mechanical-thermal processing. The minimum subgrain size decreases with decreasing recovery temperature and at 370 °C it is a b o u t 0.55 pm for the F e - 1 % V l%Ni alloys and about 0.40 pm for the F e 1%V-3%Ni alloy. (3} The addition of nickel to ferritic alloys reduces the limiting subgrain size in the

250 d e f o r m e d - a n d - r e c o v e r e d c o n d i t i o n . This observation is c o n s i s t e n t with the e n h a n c e m e n t o f cross-slip and the restriction of d y n a m i c recovery, b o t h o f w h i c h are e x p e c t e d t o o c c u r with increasing nickel c o n t e n t . (4) T h e average subgrain m i s o r i e n t a t i o n in the F e - I % V - I % N i alloy increases w i t h increasing a m o u n t s o f d e f o r m a t i o n . The measured m i s o r i e n t a t i o n s were a b o u t 1.05 ° f o r specimens d e f o r m e d t o 30% r e d u c t i o n in area, 2.1 ° for the g r o u p A specimens d e f o r m e d t o 75% r e d u c t i o n in area and 2.5 ° for the g r o u p B specimens d e f o r m e d t o 75% r e d u c t i o n in area. (5) Nickel additions e n h a n c e the t h e r m a l stability o f subgrains b y restricting subb o u n d a r y m o b i l i t y . C o n s e q u e n t l y , a smaller subgrain size is observed in the F e - l % V - 3 % N i alloys t h a n in t h e F e - l % V - 1 % N i alloys. (6) F o r materials rolled to 75% r e d u c t i o n in area, recrystallization is initiated during annealing at 4 4 0 °C and above in the F e 1 % V - l % N i alloy and at 500 °C for the F e 1 % V - 3 % N i alloy. The e l o n g a t e d subgrains serve as nuclei for recrystallization.

ACKNOWLEDGMENTS The N o r t h w e s t College and University A s s o c i a t i o n f o r Science ( N O R C U S ) and H a n f o r d Engineering and D e v e l o p m e n t L a b o r a t o r y ( H E D L ) are t h a n k e d for the financial s u p p o r t given t o o n e o f the a u t h o r s (E.R.B.) t h r o u g h the N O R C U S L a b o r a t o r y G r a d u a t e Participant P r o g r a m s p o n s o r e d b y the U.S. D e p a r t m e n t of Energy. The use o f facilities at b o t h H E D L and Battelle Pacific N o r t h w e s t L a b o r a t o r i e s was essential to this research and is greatly appreciated. REFERENCES 1 K.J. Irvine, Proc. Conf. on Strong Tough Structural Steels, Scarborough, April 1967, British Iron and Steel Research Association, London, Iron and Steel Institute, London, 1967, p. 1. 2 N.J. Petch, in B. L. Averbach, D. K. Felbeck, G. T. Hahn and D. A. Thomas (eds.), Fracture, Wiley, New York, Massachusetts Institute of Technology Press, Cambridge, MA, 1959, p. 54. 3 R. W. Armstrong, in H. Herman (ed.), Advances in Materials Research, Vol. 4, Wiley-Interscience, New York, 1970, p. 101. 4 J. D. Embury, in A. Kelley and R. Nicholson (eds.), Strengthening Methods in Crystals, Halsted, New York, 1971, p. 331.

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