Structure and deformaton behaviour of Armco iron subjected to severe plastic deformation

Pergamon PII S1359-6454(96)00156-5

Acta mater. Vol. 44, No. 12, pp. 47054712, 1996 Copyright 0 1996 Acta Metallurgica Inc. Published by Elsevier Science Ltd Printed in Great Britain. All rights reserved 1359~6454/96 $15.00 + 0.00

STRUCTURE AND DEFORMATON BEHAVIOUR OF ARMCO IRON SUBJECTED TO SEVERE PLASTIC DEFORMATION R. Z. VALIEV’,‘, YU. V. IVANISENKO*, E. F. RAUCH’ and B. BAUDELET’? ‘GEnie Physique et MCcanique des Mattriaux (URA-CNRS no. 793), Ecole Nationale SupCrieure de Physique de Grenoble, INPG, BP 46-38402 Saint Martin d’H&res Cedex, France and %stitute of Metals Superplasticity Problems, Russian Academy of Sciences, 450001 UFA Russia (Received 4 October 1995; in revised form 1 April 1996) Abstract-Structural evolutions in an Armco iron subjected to severe plastic deformation by torsion under high pressure are analysed with conventional and high resolution electron microscopes. The substructure observed at low strains appears to shrink with increasing deformation and transforms at very high strains into grain boundaries. The resulting grain size decreases down to a constant submicrometric value. Meanwhile, the material strength, as revealed by micro hardness measurements, levels out. Dislocation densities and internal stress levels are used to discuss the structural transformations. Hydrostatic pressure

and deformation temperature are believed to modify the steady-state stress level and structural size by impeding the recovery processes involving diffusion. Copyright 0 1996 Acta Metallurgica Inc.

1. INTRODUCTION The growing interest in the problem of high strains in a range of low temperatures (< 0.3 T,,,) observed in recent years is attributed to two main reasons. First, it has been established that ultra fine-grained (UFG) structures, with grain sizes as low as 20 nm, are formed by severe plastic deformation [l-4]. Materials with a structure similar to nanocrystals produced by traditional methods such as gas condensation [5] or ball milling [6], possess some rather attractive and unusual physical and mechanical properties. Secondly, in the high strain range an abnormal strain hardening has been revealed (stages IV and V [7, 81) and investigations of this nature present interest in the physics of strength and plasticity. However, deforming up to high strains at low temperature still remains a problem for most crystalline materials because of their limited ductility. Drawing or multi-stage rolling has been used to produce strains up to E = 10 [l, 91. More recently, high strains and the resulting UFG structure has been obtained by deformation modes involving shear strain such as torsion under high hydrostatic pressure [24] or equal channel angular pressing [l, 10, 111. Investigations on the material structural evolutions and strength characteristics during severe plastic deformation are of great interest since they are required for understanding both the UFG structure formation and the physical processes related to high strains. The _FTo whom

all correspondence

should

be addressed.

present points:

work

pays special

attention

to three

main

??the

structural evolution during severe plastic deformation by torsion under high pressure is analysed by transmission electron microscopy (TEM) and high resolution electron microscopy (HREM); ??hardness changes are investigated during the deformation; ??the results are compared to the data related to high strains available in the literature.

An Armco iron was taken for the investigation mainly because the structure of iron was thoroughly investigated at large drawing by Langford and Cohen [9,

121. 2. EXPERIMENTAL

PROCEDURE

Armco-iron samples, 99.95% in purity, were used for the experiments. The initial state consists of hot rolled ingots annealed at 700°C for 1 h. The resulting mean grain size is 40 pm. Washers, 0.5 mm thick and 8 mm in diameter, were then cut out from these ingots and subjected to torsion deformation under a quasi-hydrostatic pressure of P = 7 GPa at room temperature. This was achieved by compressing the samples between upper stationary and lower twisting anvils. Details on the technique have been described elsewhere [2, 4, 131. The different strain levels were achieved by turning the lower part to a definite fraction of turn. As for conventional torsion

4705

VALIEV et al.: DEFORMATION OF ARMCO IRON

4706

Fig. l(a). See caption opposite.

deformation, the disc is:

the shear strain at the outer radius R of

tAn alternative accumulated

expression, equivalent

therefore the resulting values underestimate the real deformation. Both remarks lead to the same conclusion: the strains considered in the following are y = 2nRNIL (1) only approximate and it appears more realistic to where N is the number of turns and L is the thickness refer systematically to the number of turns. of the sample. In order to compare the shear strain The samples investigated were twisted to N = l/4, with the strain obtained with other modes of l/2, 1, 5 and 10 turns. Some further specimens were added for micro hardness measurements. All samples deformation, it is convenient to convert it into the equivalent strain E,,. Using the von-Mises expression produced had a disc shape without any cracks or this gives:? cavitations. The final diameter being about 10 mm, it is seen that the equivalent strains obtained with the above equations, respectively equal to 9, 18, 36, 181 and 363, are fairly high. Two remarks should be made about equation (1): The microstructure was analysed with a JEM-2000 the first is that the deformation varies linearly from EX TEM operating at 200 kV while high resolution zero at the centre to the value quoted above for the pictures were obtained by a Hitachi-7000 HREM rim of the disc. Consequently the structures are operating at 300 kV. Thin foils were jet electroobserved at as large a radius as possible. In fact, it polished in a standard way. will be seen that for severe deformations the structure The hardness (H,) was measured on a PMT-3 looks rather homogeneous. The second point device with a pyramidal diamond indenter subjected concerns the thickness of the specimen which is to a 150 g load. Before testing the samples were observed to decrease during the test down to half its polished with 3 pm diamond paste to get a clear and initial value because of the high compressive pressure. smooth surface. No less than 13 measurements were For the sake of simplicity, the strain is calculated by made at arbitrary sites for each sample. The standard keeping the initial value of L in equation (1) and deviations did not exceed 3%. 3. EXPERIMENTAL which gives a lower value for the strain, is frequently used, i.e.

ES~= -!- 1nrJjG-i + y]. J5 It is worth recalling that, as demonstrated by Shrivastava et al. [14], this expression is not valid at large strains.

RESULTS

Let us first consider the evolutions of the structure with increasing strain as observed by TEM. Figures l(aHc) show typical structure evolutions in iron. Bright and dark field images are shown as well as diffraction pattern aspects for a constant sample area of 0.7 pm in diameter. While the structures

VALIEV et al.:

DEFORMATION

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Fig. 1. TEM micrographs of typical structures of severely deformed Armco iron: (a) N = l/4; (b) N = 1; (c) N = 5. Bright and dark field images as well as selected diffraction patterns are presented for the different states of deformation.

appear rather confused in these pictures, some important features may be inferred by comparing them. Cellular structure is typical after N = l/4 [Fig. l(a)]. The cell walls consist of tangles with a very high dislocation density close to the ultimate one for TEM resolution. Narrow walls (shown by arrows), typical for subgrain structure, are also observed but only in a few cases. Electron diffraction patterns of this state are typical for a cell structure with weak misorientations

between neighbouring volumes. The mean cell size as determined by dark field images is 360 nm. For one half and one turn, the microstructt tres are refined: the cell size decreases down to 2130 and 210 nm respectively while the number of ‘n arrow’, sharp boundaries increases. Misorientations ai -e more evident, and are especially prominent when IV= 1, for which the diffraction pattern is similar to that of a fine-grained structure [Fig. l(b)].

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After five turns an ultrafine-grained structure is observed with a mean grain size of about 100 nm [Fig. l(c)]. Of importance is the fact that no change upon the structural appearance was noted for further increases of straining up to N = 10. Surprisingly, for such a level of strain, dislocations are hard to observe inside the small grains even at higher magnifications [Fig. 2(a, b)]. Rather, highly dispersed defects with a circular shape were revealed. These defects are believed to be dislocation debris and indeed they disappear after grain growth onset when the samples are heated. Similar defects in strongly deformed metals have been commented on in the literature [15]. The structure is also characterized by high internal stresses, as indicated by bent extinction contours inside grains. High internal stresses and microdistortions were also measured in ultra fine-grained Cu and N&Al processed by severe plastic deformation [3, 161. Specific features of the given structural state were investigated by the HREM method. The data on structures at and near grain boundaries obtained by direct resolution of lattice fringes are of special interest. Iron has a b.c.c. structure with a lattice spacing of 0.286 nm. Only (110) planes with the interface distance d = 0.203 nm could be resolved in the HREM operating at 300 kV (C, = 0.9 mm). Figure 3(a) shows an example of a highly distorted region situated near a grain boundary in iron deformed to 10 turns. The values of elastic strains, as determined from the lattice fringe bending, reach 3 to 5%. Recently similar values of elastic strains have been estimated by the analysis of Moire patterns during the HREM studies of UFG nickel prepared by severe plastic deformation [17]. As known, the presence of dislocations may be indicated by HREM through the direct observation of crystal lattice extraplanes. Unfortunately, the typical b.c.c. dislocations with { 111) Burgers vectors are not revealed by imaging only one family of (110) planes, as is performed on the pictures. A high dislocation density will rather give rise to a very complex, amorphouslike contrast of the lattice fringes and indeed, such types of observations are frequently made on these materials [Fig. 3(b)]. The fact that dislocation densities up to 10” mm2 were observed at and near grain boundaries in a f.c.c. material (Al-3%Mg, [IS]) deformed in similar conditions leads us to the conclusion that the typical aspect shown in Fig. 3(b) is indicative of dislocation concentrations. Let us consider the hardness measurements from deformed samples. The H, versus N dependence is shown in Fig. 4. The hardness in the less deformed sample is equal to 3.25 GPa and grows with increasing strain up to N = 5 where it reaches a constant value of about 4.6 GPa. For N > 5, the values of H, measured along the diameter of the sample are rather homogeneous. This fact confirms that the strength of the material becomes strain independent at large deformations.

OF ARMCO

IRON

(4

@I Fig. 2. TEM micrographs with large magnifications illustrating: (a) the absence of visible dislocations inside grains-bright field image. (b) The presence of small debris (arrowed) in the structure in highly deformed iron (N = IO)-dark field image.

4. DISCUSSION

The structural evolutions in severely deformed iron were reported by Langford and Cohen [9, 121 for room temperature drawing. Mainly, they noted a continuous decrease of the transverse cell size accompanied by a monotonous increase of the flow

VALIEV et al.:

DEFORMATION

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Fig. 3. HRI 3M images of lattice fringes near grain boundaries in highly deformed iron (N = 10) (a) Elastic crystal lattice distortions near a grain boundary. Note the deviations between the (110) planes and the line drawn; (b) another area where the complex of lattice plane images is seen (A).

stress. No saturation was observed in their work. The main difference between their deformation mode and the present one lies in the strain levels: while drawing was performed until an equivalent strain of 10, the deformation calculated with equations (1) and (2), although approximate, is considerably higher. Moreover, the strain path related to drawing differs from the simple shear employed in the present work and therefore specific material and crystallographic rotations are expected for each case. As a consequence, the steady state observed after torsion under hydrostatic pressure is not in contradiction with their findings. It is worth emphasizing that constant strength was reported for f.c.c. metals at

room temperature [2, 191, but to the authors’ knowledge, for b.c.c. materials there is only evidence that saturation may be reached [20]: this also concerns torsion on an iron base steel but without pressure, and the steady state is attained at E = 5 while the stress level is about 0.7 GPa. From the present results it appears that the stress saturation is correlated to structural constancy and of importance is the fact that, here, the microstructure reaches a steady state only after grain refining. Let us analyze in detail the relation between structural evolutions and strength changes during intense plastic deformation for iron. For this purpose, the usual approximate relation between

VALIEV et al.: DEFORMATION OF ARMCO IRON

4710

5.0

Estimated 87 174

0

cl F5 4.5.

2

348

435

*

*

*

1

strain 251

* 4.0

*

F Ts 3.5 2 * 3 00

*

*

2

4 Number

6

8

of

turns

10

12

Fig. 4. Vickers hardness as function of number of turns. The corresponding deformation values are added.

microhardness and flow stress, namely H, z 30~,, will be employed. The contribution of the grain boundaries to the stress level may be estimated with the Hall-Petch relationship: a, = co + k&0.’ where d is the grain size. Because the Armco iron displays a yield point behaviour, there are some doubts about the correct values to be chosen for e. and k. Using the values determined at yield point [21], i.e. co x 76 MPa and 15.8 < k < 17.7 N mm-“* respectively, one obtains 1.65 < cry< 1.85 GPa at N = 5 turns (d = 100 nm). The corresponding hardness is higher but of the same order of magnitude as the experimental one (HvcXp = 4.7 GPa, a, = 1.57 GPa). By contrast, the measurements performed behind the yield point (e.g. a0 z 300 MPa and k = 12.6 N mm-3’2, [22]) lead to a stress level even closer to the measured ones (oY= 1.56 GPa) with, however, some difficulties for separating the grain size from the dislocation strengthening effect. The above values lead to the conclusion that the high strength of ultrafine-grained iron may be related to the grain size. By contrast, at intermediate levels of strain, an intragranular hardening is expected. Indeed, following the TEM observations, cells rather than grains are obtained for N I 1. The substructure

Fig. 5. Schematic

strengthening follows the relation a, = go’+ k’d-’ [9,23] and for iron the parameters of this expression are aA = 34 MPa and k’ = 0.122 N mm-’ [9]. A mean cell size of 360 nm (N = l/4) gives uj = 0.37 GPa. The calculated hardness (H, = 1.1 GPa) is three times less than the recorded value H,.,, = 3.25 GPa, a discrepancy too high to be entirely attributed to the imprecise hardness/flow stress relation. The Hall-Petch relation for this case leads to 2.7 < H, < 3 GPa, a value which again has a good order of magnitude. This hints at the idea that the structure contains not only cell walls but also some amount of highly misorientated boundaries. The structure evolution observed in iron is in good agreement with modern concepts on structural changes occurring in metals during large plastic deformation. The tendency to a decrease of cell size and wall widths, as well as an increase of misorientations was noted in a lot of works devoted to high strains in low temperature ranges [7, 15, 24,251. Some experimental works give evidences on the presence of high angle misorientations in strongly deformed structures [7, 25, 261. The ultra-fine grains observed in our case appear to be the result of increasing misorientations between cells leading to a granular type structure. However, the grain boundaries that were produced during deformation differ somewhat from those obtained at recrystallization. Mainly, high distortions and high internal stresses are observed in their neighbourhood. The experimental facts suggest the following scheme for structural evolutions during intense plastic straining of iron (Fig. 5). Large strains lead to an increase in the number of dislocations that are known to concentrate in cell walls. The decrease of the dislocation wail thickness is usually attributed to a recovery process which permits segment rearrangements. As annihilation of dislocations of opposite sign is possible, only excess dislocations of mainly one sign (for each active system) are stored at cell boundaries. These boundaries are probably not as regular as usual grain boundaries (and therefore cause the growth of internal stresses) but rather are composed of accumulated dislocations gathered in a

model for the evolution of dislocation structures at different deformation intense plastic straining in Armco iron (see text for details).

stages during

VALIEV

et al.:

DEFORMATION

small interface volume. The increasing resultant misorientation is believed to be responsible for the observed granular type structure. However, the storage of excess dislocations and consequently the refinement of the structure is obviously limited by another recovery process. Indeed, the dislocation density in the interface volume is expected to be bounded by some critical value. When dislocations of opposite signs are considered, Essmann and Mughrabi [27] consider that 1014-1017mm* are values for such a (static) critical density in copper. For the excess dislocations, it may be imagined that spontaneous rearrangement of dislocations with different Burgers vectors also occurs when the distance between them becomes small enough. Further experimental evidence is required to determine the nature of the rearrangement mechanism. One possible process would imply dislocation motion along grain boundaries and preferential annihilation at triple junctions leading to grain boundary sliding. It may also be considered that dislocation cores are no more well defined for high densities and that atomic rearrangments are involved in the final steps of deformation. In any case, when the rate of dislocation accumulation and absorption on grain boundaries is similar, a strain independent structure should arise. This occurs for N larger than five turns where a so-called ‘dissipative’ structure is obtained and the microhardness value levels out. The above scheme is in good agreement with a recent investigation of the mechanical behaviour of UFG Cu produced by severe plastic deformation [28]. The latter has also revealed flow stress saturation at very large strains and has shown some evidence of grain boundary sliding at rather low temperature (0.21 T,). It is well known that the stress saturation is a typical phenomenon for hot or even warm deformations when dynamic recrystallization or easy recovery processes take place. As mentioned above, the steady state is sometimes observed at room temperature and occurs after specific steps on the hardening curve (stages IV and V in f.c.c. metals [7, 81). It is worth emphasizing that, in all these works, the structure is stable because of rearrangement processes occurring in cell walls while in the present case the steady state is related to the formation of an ultra fine granular structure. Obviously, the latter is obtained only when the recovery processes active in the other cases (which lead to the cellular structure) are not sufficiently efficient to promote the steady state. The problem is to understand why this is so for the present mode of deformation. A tentative explanation is that the maximum stress that can be reached for a given metal is sensitive to the hydrostatic pressure. As known, the applied pressure can considerably slow down diffusion and, consequently, delay recovery kinetics. Let us estimate a possible retardation of diffusion in our case through the following relationship between

OF ARMCO

the self-diffusion pressure P [29]:

IRON

activation

4711

energy

Q and

the

Q = Qo + uPb3, where b3 is the atomic volume, and CLis a coefficient dependent on the crystal lattice type. Taking b = 0.248 nm, a = 0.6 [29] and P = 7 GPa for iron we find that the activation energy of self-diffusion changes by 30%, i.e. a considerable delay in recovery processes is expected during the test. This can explain the smaller grain size and the higher stress level as compared with that obtained by torsion (without pressure). Also of interest is the fact that for strains less than 10, conventional torsion gives rise to a steady state while drawing does not. Indeed, the latter is expected to promote a pressure of the order of magnitude of the flow stress (around 1 to 2 GPa for iron) while torsion is not.

5. CONCLUSIONS

(1) A strain insensitive microhardness

indicating Bow stress saturation has been recorded for Armco iron samples subjected to severe plastic deformation by torsion under 7 GPa hydrostatic pressure at room temperature. (2) TEM and HREM observations have shown a progressive evolution of the cellular structure to an ultra fine-grained one in the course of straining. The structure size (i.e. initially the cell diameter and then the grain size) is found to decrease down to a value that does not change during further straining as soon as the stress saturation is attained. (3) The grain interiors show some evidence of high internal stresses which are attributed to high dislocation densities at and near grain boundaries. of very high strains to (4) The requirement promote a steady state is believed to be due to the delaying effect of low temperature and high hydrostatic pressure on diffusion assisted recovery processes. Acknowledgement-One of the authors (RZV) would like to thank Dr W.-A. Chiou for the help with the HREM studied during his work on COBASE program at Professor J. P. Weertman’s laboratory at North Western University, Evanston, U.S.A.

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