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Texture evolution during recrystallization in a boron-doped Ni76Al24 alloy
Materials Science and Engineering A277 (2000) 1 – 10 www.elsevier.com/locate/msea
Texture evolution during recrystallization in a boron-doped Ni76Al24 alloy Sandip Ghosh Chowdhury a,*, R.K. Ray b , A.K. Jena b b
a Materials Characterization Di6ision, National Metallurgical Laboratory, Jamshedpur, 831007, India Department of Materials and Metallurgical Engineering, Indian Institute of Technology, Kanpur, 208016, India
Received 27 July 1999; received in revised form 6 August 1999
Abstract The evolution of recrystallization texture in the intermetallic compound Ni76Al24(B) has been investigated using the conventional pole figure as well as orientation distribution function (ODF) methods. The initial L12 structure which transforms to DO22 on cold rolling reverts back to the L12 during annealing. The annealing process can be divided into three stages: recovery, reordering and recrystallization. The moderately strong deformation texture resulting from cold rolling becomes very weak during the recovery and reordering processes by an ‘in-situ’ orientation change due to reordering that precedes recrystallization. The texture remains weak throughout recrystallization and grain growth stages. TEM investigations show evidence of twin like features during the reverse DO22 L12 transformation. This is expected to lead to fragmentation of the matrix grains resulting in a weak texture. The prominent recrystallization texture components are {025} B 100\ and {011}B 100\ , of which the latter component survives best with continued annealing. © 2000 Elsevier Science S.A. All rights reserved. Keywords: Texture; Recrystallization; Boron-doped Ni76Al24 alloy
1. Introduction In recent years, boron-doped polycrystalline Ni3Al has evoked considerable interest for high temperature structural applications. However, the deformation as well as recrystallization behaviour of this material and the textural changes associated with them have not received much attention yet due to the lack of proper workable material. It is now known that by controlling the grain size to the order of few microns, through proper thermomechanical treatments, the material could be cold rolled without noticeable crack formation [1,2]. In a previous communication [3], the structural and textural changes during cold rolling of a B-doped Ni76Al24 alloy were reported. The present paper describes the recrystallization behaviour of the same alloy, previously subjected to heavy amounts of cold rolling. Previous work suggests [3] that the phenomenon of recrystallization in ordered intermetallics of the type of * Corresponding author. Tel.: +91-657-426091; fax: + 91-657426527/4312. E-mail address: [email protected] (S. Ghosh Chowdhury)
Ni76Al24(B) follows the usual sigmoidal behaviour. However, the rate of recrystallization in such materials is expected to be quite slow. The long range order present in such materials is bound to have a strong influence on both the mechanism and the kinetics of recrystallization. All these aspects, therefore, should be looked into quite carefully for a proper understanding of the phenomenon. Gottstein et al. [1] showed that the textural components determined by them during recrystallization of a heavily cold rolled Ni76Al24(B) were dependent on the annealing temperature. They attributed this unusual behaviour to low grain boundary mobility. Later, Ball and Gottstein [4] also observed that the recrystallization texture is very weak, almost of random intensity. They suggested this to be due to a high nucleation rate, low grain boundary mobility and annealing twin formation in this material. In a later publication, Ball and Gottstein [5] reported partial disordering of the alloy by cold rolling and proposed that microband formation could be the main cause for the measured disordering. However, they did not mention anything regarding changes in the long range order parameter during an-
0921-5093/00/$ - see front matter © 2000 Elsevier Science S.A. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 9 9 ) 0 0 5 6 4 - X
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S. Ghosh Chowdhury et al. / Materials Science and Engineering A277 (2000) 1–10
Fig. 1. X-ray line profiles of boron-doped Ni76Al24 as function of degree of cold rolling.
nealing; neither did they make a conscious effort to study the effects of long range order on texture development. The present work was undertaken with a view to filling in the gaps in the existing literature on the recrystallization behaviour of B-doped polycrystalline Ni3Al alloys.
free surface of the flat samples were chosen for this purpose. The annealed microstructures were characterized using a JEOL 2000 FX II transmission electron microscope (TEM) operated at 160 kV. The variation, during annealing, of order parameter (S) was monitored by X-ray diffraction technique. This parameter (S) was calculated from the integrated intensity ratios of the (100)/(200) and (110)/(220) pairs of reflections. The relationship used between S and the above intensity ratios is as follows [6]:
2. Experimental procedure The material used was the intermetallic alloy Ni76Al24 with 0.24 at.% boron. The arc melted alloy in the cold forged condition was obtained (courtesy Professor Dr. G. Gottstein of the Institut fu¨r Metallkunde und Metallphysik, RWTH Aachen, Germany). After homogenization at 1050°C for 25 min, the material yielded an average grain size of 30 mm. Homogenized samples of 4.5-mm thickness were cold rolled at room temperature to a reduction of 85%. Isochronal annealing of the cold rolled material was carried out from 100 to 1000°C at intervals of 100°C. The holding time was 1 h in each case. A few samples were also isothermally annealed at four different temperatures, namely 750, 800, 850 and 900°C for various lengths of time. Thin foils were prepared from the materials at different stages of annealing by mechanical polishing, followed by electropolishing in a jet electropolisher using a solution of 10% H2SO4 in methanol. Rolling plane sections at a depth of 1/4 the thickness from the outer
Fig. 2. Variation of the order parameter, S, with degree of cold rolling.
S. Ghosh Chowdhury et al. / Materials Science and Engineering A277 (2000) 1–10
Fig. 3. Variation of order parameter, S, during 1-h isochronal annealing of the 85% cold rolled material.
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subscripts s and f refer to the superlattice and fundamental reflections, respectively. The rationale behind the choice of the (100)/(200) and (110)/(220) line pairs is that because the planes are parallel in any such pair, the value of the intensity ratio will not be affected by textural changes (although the absolute intensities of the lines will be obviously influenced by texture). For texture measurement, samples of 14×24 mm2 were cut out from the flat plates of different annealed materials. From each specimen, material was etched off from one side up to 1/4 of the thickness, and texture was measured on that surface only. In each case, a set of four pole figures, such as {111}, {200}, {220} and {311}, were measured and the orientation distribution function (ODF) was computed by means of series expansion [7] to an order of l= 22 and plotted in the f2 = constant sections.
3. Results and analysis
3.1. Structural changes
Fig. 4. Variation of the order parameter, S, with time at the isothermal annealing temperature of 850°C.
S2=
Is (xAl fAl +xNi fNi)2 (Lp)f (e − 2M)f 16 If ( fAl −fNi)2 (Lp)s (e − 2M)s
where, I is the measured integrated intensity; xAl and xNi are the atom fractions of Al and Ni, respectively, in the alloy; fAl and fNi are the atomic scattering factors for Al and Ni, respectively; Lp is the Lorentz-Polarization factor; and e − 2M is the Debye-Waller factor. The
3.1.1. Order parameter Fig. 1 shows the X-ray diffraction profile taken from the alloy, in the initial homogenized condition and after different degrees of cold rolling. The pattern for the homogenized alloy has been analyzed as that for an L12 structure, showing clear (100) and (110) superlattice peaks over and above the fundamental lines. With an increase in the degree of cold rolling, the (100) peak becomes weaker. The order parameter measured from (100)/(200) intensity ratio reduces to zero after 85% cold work. However, the intensity of the other superlattice peak (110) persists even after 85% cold rolling. In fact, the order parameter calculated from (110)/(220) intensity ratio gradually decreases and assumes a nearly constant value of 0.45 from 65% deformation onwards. The variations of the order parameter with amount of
Fig. 5. Variation of hardness with time during isothermal annealing at four different temperatures. Arrows indicate the start of recrystallization.
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The variations in the order parameter values during the course of 1-h isochronal annealing of the 85% reduced material are depicted in Fig. 3. It is observed that the intensity of the (100) peak remains zero even after isochronal annealing at 200°C. With the increase in temperature the value of the order parameter, S, measured from the (100)/(200) intensity ratio, increases rapidly until 400°C and then shows a gradual increase and attains the maximum value of 0.8 at 1000°C (Fig. 3). By contrast, the value of S from the (110)/(220) line pair shows a gradual increase with temperature to the initial value (before cold rolling) of 0.8. The above isochronal annealing curves clearly indicate a major change in order parameter between 200 and 400°C in the 85% cold rolled sample. The variation of S during isothermal annealing is shown typically in Fig. 4 for the sample annealed at 850°C. The variation of hardness during isothermal annealing at four different temperatures as depicted in Fig. 5 clearly shows that nearly 66% of the hardness recovers prior to the beginning of the recrystallization process.
3.1.2. Microstructure The cold rolled (85%) microstructure of this alloy was characterized by a high density of shear bands as clearly shown in typical optical micrographs (Fig. 6a). TEM micrographs have indicated the presence of a high density of dislocations, microbands as well as deformation twins (Fig. 6b). The SAD pattern from the twinned area in Fig. 6b is shown in Fig. 6c. Annealing at 500°C shows at many places groups of parallel stripes running nearly perpendicular to original grain boundaries (Fig. 7). Features similar to these stripes were observed in a Cu3Pt alloy during transformation from the disordered to the L12 structure [9]. Those stripes were identified by the authors as ordered domains of the L12 structure. It has already been seen from X-ray diffraction results (Fig. 3) that at this
Fig. 6. (a) Optical micrograph showing shear bands. (b) TEM micrograph showing high densities of dislocations, microbands and deformation twins. (c) SAD pattern of twinned area in (b).
rolling reduction have been plotted in Fig. 2. The above behaviour has been found to be consistent with a structural transformation from L12 to DO22 taking place in the material due to rolling deformation [8]. It might be mentioned that peak broadening due to shrinking crystallite size and increasing RMS strain during cold work is typical of essentially all metals. This fact may tend to ‘blur’ the superlattice peaks and hence, exaggerate the disordering. Perhaps the lower modulus value along B 100 \ contributes to the distinction between the two families of planes.
Fig. 7. TEM micrograph of stripe like features emanating from the grain boundaries after annealing at 500°C (Moire´ fringes like features near a grain boundary).
S. Ghosh Chowdhury et al. / Materials Science and Engineering A277 (2000) 1–10
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Fig. 9. TEM micrograph of a group of grains after annealing at 700°C.
Fig. 10. TEM micrograph of a partly recrystallized area after annealing at 800°C.
Fig. 8. (a) TEM micrograph of twin like features after annealing at 500°C, (b) SAD pattern of the same area and (c) key to the SAD pattern.
temperature, the (100) superlattice peak which disappeared after heavy rolling deformation, reappears and the structure becomes L12 again. Presumably, annealing leads to a retransformation from the DO22 to the L12 structure. A distinct Moire´ fringe like feature sometimes observed at grain boundary regions (shown with an arrow in Fig. 7) can be speculated to indicate the possible coexistence of the above two structures. Twin like features are also sometimes observed in regions which have recovered sufficiently. This is shown typically in Fig. 8a. The SAD taken from this area (Fig. 8b) shows the [011] pattern of single crystal of the L12
Fig. 11. f2 sections of the ODF of 85% cold rolled boron-doped Ni76Al24.
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Fig. 12. f2 sections of the ODF after annealing at 850°C for 17 h.
tallization aids in the formation and growth of the Goss oriented regions. The ODF of the alloy, annealed at 1025°C for 1 h, again shows a Goss component along with rotated cube orientations (Fig. 13c). The texture components reported by different investigators (including the present work) during recrystallization of heavily rolled Ni76Al24(B) have been tabulated in Table 1. The variation of volume fraction recrystallized with time at various annealing temperatures is shown in Fig. 14 for the 85% cold rolled material [10]. As shown in Fig. 14 the alloy had yet to start recrystallizing after 1 h of annealing; only partial recrystallization set in after annealing for 5 and 11 h, whereas 24-h annealing produced slight grain growth beyond completion of recrystallization. In order to understand the effects of the structural change and the change in order during annealing on texture, measurements were made on samples isothermally annealed at 850°C for 1, 5, 11 and 24 h. The
structure but does not show distinct twin spots. The key to the diffraction pattern is given in Fig. 8c. Annealing at 700°C leads to copious nucleation of large number of fine grains (Fig. 9). Fig. 10 again shows a typical partly recrystallized area showing recrystallization twins inside. The SAD from the central grain shows the [011] zone axes of the L12 structure.
3.2. De6elopment of texture Texture measurements were undertaken on the partly and fully recrystallized samples which were initially cold rolled by 85%. As reported earlier [3], the texture of the 85% cold rolled material (Fig. 11) was reasonably sharp (maximum intensity 5 R). The texture after primary recrystallization has been found to be invariably weaker (maximum intensity 1.5 R). Fig. 12 shows the ODF of a fully recrystallized (metallographically) sample, annealed at 850°C for 1 h. A comparison between the above two figures clearly indicates that the texture severity drastically decreases upon recrystallization and that the recrystallization texture of this material is very weak. Even then some distinct features are discernible in the ODF of the recrystallized sample, such as the presence of a rotated cube {025}B 100\ along with a weak Goss {011} B100 \ component. The ODF of the sample recrystallized at 950°C for 1 h, however, shows only the weak rotated cube orientation, but no Goss orientation (Fig. 13a). The Goss orientation develops a bit later in this material, after 48 h of annealing (Fig. 13b) and there is a considerable amount of scatter around this orientation. The intensity of this orientation in this sample is also markedly higher than its intensities in the ODFs of the earlier samples. This clearly indicates that grain growth after primary recrys-
Fig. 13. (a) f2 =0° sections of the ODF after annealing at 950°C for 1 h. (b) f2 =0° sections of the ODF after annealing at 950°C for 48 h. (c) f2 =0° sections of the ODF after annealing at 1025°C for 1 h.
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Table 1 Effect of temperature of annealing on the development of recrystallization texture components in Ni76Al24(B) Material
Temperature (°C)
Time
Component
Intensity recrystallization/roll
Reference
90% Cold rolled
650 700 750 710 750 850 950 950 1025
1 1 1 90 30 17 1 48 1
{310}B135\ {211}B135\ {012}B121\, as above {013}B100\, {025}B100\, {025}B100\ {011}B100\ {025}B100\,
3.3/16.2 3.8/16.2 2.0/6.3 3.4/6.3 2.0/6.8 1.6/5.0 1.5/5.0 1.8/5.0 1.6/5.0
[1] [1] [4] [4] [4] Present Present Present Present
92% 90% 92% 85%
Cold Cold Cold Cold
rolled rolled rolled rolled
h h h min min h h h h
{112}B131\ {102}B201\, {112}B294\ {011}B100\
{011}B100\
work work work work
Fig. 14. Variation of the volume fraction recrystallized with time at various annealing temperatures.
f2 = 0° sections of the ODFs of these four samples are presented in Fig. 15a – d. All these ODFs are found to be essentially similar in the sense that all are very weak, much weaker than the ODF of the 85% cold rolled material (Fig. 11). The reproducible textural components after 1 h of annealing are {025} B 100 \ and {102}B 201\. After 5 h, a weak (near) Goss component {011}B 311\ is also observed in addition to the above components. After 11 h of annealing, all the previous components are obtained, but here the (near) Goss component {011}B 311 \ is found to be the most intense and has the exact Goss component also in its spread. The {025}B100 \ component is non-existent after 24-h annealing, indicating that the (near) Goss component of texture survives best on continued annealing. The development of recrystallization texture in boron-doped Ni76Al24 alloy is very different from that of pure Ni which is known to exhibit a strong cube {100}B 001\ texture under identical conditions. Both the materials have nearly the same stacking fault energy (SFE) as mentioned in Refs. [1] [2]. Irrespective of that, the deformation as well as recrystallization textures have been found to be widely different in the two cases. Presumably, factors such as ordering of Ni76Al24(B) and
the structural changes it undergoes during deformation (L12 DO22) as well as during annealing (DO22 L12) far outweigh the effect of SFE in determining the final textures. The 1-h annealed sample represents only a recovery stage. However, the S parameter was found to
Fig. 15. f2 =0° sections of the ODFs after annealing at 850°C: (a) 1 h, (b) 5 h, (c) 11 h and (d) 24 h.
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S. Ghosh Chowdhury et al. / Materials Science and Engineering A277 (2000) 1–10
Fig. 16. XRD line profiles for samples cold rolled 85% and then isothermally annealed at 850°C for 5 and 30 min.
Fig. 17. DSC thermogram of 85% cold rolled boron-doped Ni76Al24.
increase from zero (after 85% cold rolling) to a value of 0.64 at this stage (Fig. 4). The XRD line profile for the samples annealed at 850°C for 5 and 30 min as well as for the 85% cold rolled sample have been plotted in Fig. 16. Clearly, a structural reordering from the DO22 to L12 structure has taken place during the above
annealing stages, which do not however, show any perceptible recrystallization. Thus, in a way, the formation of the weak annealing texture from a strong deformation texture can be termed as a kind of ‘transformation’ texture produced by the structural change from DO22 back to L12 during annealing.
S. Ghosh Chowdhury et al. / Materials Science and Engineering A277 (2000) 1–10
4. Discussion On the basis of what has been stated above, the annealing process in this alloy appears to involve a combination of three different phenomena: recovery, reordering and recrystallization. Several aspects need to be considered in explaining the results of the recrystallization texture measurements. The first pertains to the fact that a very weak texture is produced from the sharp cold rolling texture during the ‘recovery’ stage, which does not involve any reorientation of grains as a whole, but during which substantial atomic shuffling takes place producing the parent L12 structure from the DO22, which was the end structure after cold rolling. It is known that the DO22 structure can be derived from the L12 by the introduction of a 1/2[110] APB on every {001} plane. Similarly, the introduction of an APB on every (001) plane in the DO22 structure again transforms it back into the L12 structure [11]. The drastic weakening of the texture intensity after 1-h annealing at 850°C [ f(g)= 1.7] from that of the 85% cold rolled alloy [ f(g)= 5.0] can not be related to any reorientation of grains (since no recrystallization takes place at this stage), but, on the other hand, could be related to the DO22 L12 reverse transformation. It may also be recalled that annealing has been found to be associated with the appearance of a large density of twin like features. Tanner and Ashby [12,13] made a theoretical analysis of the relief of strains associated with certain phase transformations in intermetallic alloys, such as Ni3V which has a DO22 structure at room temperature. According to them, when ordering in a binary alloy involves a change in crystal symmetry the reaction is invariably accompanied by an accumulation of substantial internal strains. Under certain conditions, these strains can be relieved through twinning of the product phase. Thus, the twin like features obtained in the product phase after annealing (long before recrystallization takes place) in the present boron-doped Ni76Al24 alloy (Fig. 8a) can be assumed to be due to such an effect. It has been shown [14] that during annealing, the heat evolution before the onset of recrystallization is quite substantial (11.5 kJ/mol) (peak A in Fig. 17). In comparison to that, the energy expended during recovery for Ni varies from 3.7 to 5.4 kJ/mol [15]. Therefore, a major part of the stored energy release in the present material can be attributed to the DO22 L12 reverse transformation. This clearly indicates that significant relief of strain is associated with this transformation, and it is no wonder, therefore, that a high density of twin like features accompanies this transformation. The incidence of twinning will further lead to fragmentation of the grains, resulting in a weakening of the texture intensity after 1 h an.
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nealing at 850°C from that of the 85% cold rolled alloy. The weak texture obtained after annealing for 1 h at 850°C continues as it is during progressive annealing up to 5, 11, 17 and 24 h, i.e. during the stages of partial recrystallization to full recrystallization, continuing even up to the grain growth stage. Very weak recrystallization texture in Ni3Al(B) has also been reported by Ball and Gottstein [2]. The almost random recrystallization texture in this alloy has been attributed by them to the high nucleation rate and the low grain boundary mobility, which produces a lack of growth selection during recrystallization. Using this argument it is difficult to reason out why the recrystallization texture should be so very weak compared to the cold rolling texture and at the same time why it should not bear any resemblance at all with the essential features of the cold rolling texture. On the other hand, the results of the present investigation suggest that when recrystallization starts during annealing, the sharp cold rolling texture has already been replaced by a weak ‘transformation’ texture. Subsequent recrystallization taking place in such a weakly textured matrix is expected to give rise to only a weak recrystallization texture. Of all the orientations observed in the ODFs of the alloy during partly or fully recrystallized conditions, the Goss component {011}B 100\ appears to be the most prominent. It is known that this component usually originates in deformation/shear bands. In the present alloy, formation of shear bands started from 65% cold rolling onwards and the density of such bands increased significantly after 85% cold work [3].
5. Conclusions The DO22 structure (the end product after 85% cold rolling) reverts back to the initial L12 during annealing. The reordering process is accompanied by the formation of twins which might accommodate the transformation strain. During annealing the texture changes drastically even before the onset of recrystallization and becomes very weak, almost random. This can be correlated with the DO22 L12 transformation. The formation of twins during the transformation leads to fragmentation of the matrix resulting in a weak texture. During recrystallization a weak recrystallization texture is produced from the already recovered and reordered material which itself has a weak texture. The prominent recrystallization texture components are {025}B100\ and {011}B 100\ . The Goss component becomes prominent in the post-recrystallization stage.
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Acknowledgements The authors are indebted to Professor G. Gottstein, Director, Institute fu¨r Metallkunde und Metallphysik, RWTH Aachen, Germany for the provision of laboratory facilities for conducting the texture work. Acknowledgements are also due to V. Kumar and Mr. Umashankar for their help during the experimental work.
References [1] G. Gottstein, P. Nagpal, W. Kim, Mater. Sci. Eng. A108 (1989) 165. [2] J. Ball, G. Gottstein, Intermetallics 1 (1993) 171. [3] S. Ghosh Chowdhury, R.K. Ray, A.K. Jena, Mater. Sci. Eng. A A246 (1998) 289.
.
[4] J. Ball, G. Gottstein, Intermetallics 1 (1993) 191. [5] J. Ball, G. Gottstein, Intermetallics 2 (1994) 205. [6] B.D. Cullity, Elements of X-Ray Diffraction, Narosa, London, 1993. [7] H.J. Bunge, Texture Analysis in Materials Science, Butterworth, London, 1982. [8] S. Ghosh Chowdhury, R.K. Ray, A.K. Jena, Scripta Metall. Mater. 32 (1995) 1501. [9] K. Mitsui, Y. Mishima, T. Suzuki, Phil. Mag. A53 (1986) 357. [10] S. Ghosh Chowdhury, A.K. Jena, R.K. Ray, Metall. Mater. Trans. A 29A (1998) 2893. [11] M. Yamaguchi, Y. Umakoshi, in: R.W. Cahn, P. Hassen, E.J. Kramer (Eds.), Plastic Deformation and Fracture of Solids, VH, Weinheim, 1993, p. 251. [12] L.E. Tanner, Phys. Stat. Sol.(a) 30 (1968) 685. [13] L.E. Tanner, M.F. Ashby, Phys. Stat. Sol. 33 (1969) 59. [14] S. Ghosh Chowdhury, Ph.D Thesis, I.I.T. Kanpur, India, 1995. [15] L.M. Clarebrough, M.E. Hargrieves, G.W. West, Proc. R. Soc. A232 (1955) 252.
Texture evolution during recrystallization in a boron-doped Ni76Al24 alloy Sandip Ghosh Chowdhury a,*, R.K. Ray b , A.K. Jena b b
a Materials Characterization Di6ision, National Metallurgical Laboratory, Jamshedpur, 831007, India Department of Materials and Metallurgical Engineering, Indian Institute of Technology, Kanpur, 208016, India
Received 27 July 1999; received in revised form 6 August 1999
Abstract The evolution of recrystallization texture in the intermetallic compound Ni76Al24(B) has been investigated using the conventional pole figure as well as orientation distribution function (ODF) methods. The initial L12 structure which transforms to DO22 on cold rolling reverts back to the L12 during annealing. The annealing process can be divided into three stages: recovery, reordering and recrystallization. The moderately strong deformation texture resulting from cold rolling becomes very weak during the recovery and reordering processes by an ‘in-situ’ orientation change due to reordering that precedes recrystallization. The texture remains weak throughout recrystallization and grain growth stages. TEM investigations show evidence of twin like features during the reverse DO22 L12 transformation. This is expected to lead to fragmentation of the matrix grains resulting in a weak texture. The prominent recrystallization texture components are {025} B 100\ and {011}B 100\ , of which the latter component survives best with continued annealing. © 2000 Elsevier Science S.A. All rights reserved. Keywords: Texture; Recrystallization; Boron-doped Ni76Al24 alloy
1. Introduction In recent years, boron-doped polycrystalline Ni3Al has evoked considerable interest for high temperature structural applications. However, the deformation as well as recrystallization behaviour of this material and the textural changes associated with them have not received much attention yet due to the lack of proper workable material. It is now known that by controlling the grain size to the order of few microns, through proper thermomechanical treatments, the material could be cold rolled without noticeable crack formation [1,2]. In a previous communication [3], the structural and textural changes during cold rolling of a B-doped Ni76Al24 alloy were reported. The present paper describes the recrystallization behaviour of the same alloy, previously subjected to heavy amounts of cold rolling. Previous work suggests [3] that the phenomenon of recrystallization in ordered intermetallics of the type of * Corresponding author. Tel.: +91-657-426091; fax: + 91-657426527/4312. E-mail address: [email protected] (S. Ghosh Chowdhury)
Ni76Al24(B) follows the usual sigmoidal behaviour. However, the rate of recrystallization in such materials is expected to be quite slow. The long range order present in such materials is bound to have a strong influence on both the mechanism and the kinetics of recrystallization. All these aspects, therefore, should be looked into quite carefully for a proper understanding of the phenomenon. Gottstein et al. [1] showed that the textural components determined by them during recrystallization of a heavily cold rolled Ni76Al24(B) were dependent on the annealing temperature. They attributed this unusual behaviour to low grain boundary mobility. Later, Ball and Gottstein [4] also observed that the recrystallization texture is very weak, almost of random intensity. They suggested this to be due to a high nucleation rate, low grain boundary mobility and annealing twin formation in this material. In a later publication, Ball and Gottstein [5] reported partial disordering of the alloy by cold rolling and proposed that microband formation could be the main cause for the measured disordering. However, they did not mention anything regarding changes in the long range order parameter during an-
0921-5093/00/$ - see front matter © 2000 Elsevier Science S.A. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 9 9 ) 0 0 5 6 4 - X
2
S. Ghosh Chowdhury et al. / Materials Science and Engineering A277 (2000) 1–10
Fig. 1. X-ray line profiles of boron-doped Ni76Al24 as function of degree of cold rolling.
nealing; neither did they make a conscious effort to study the effects of long range order on texture development. The present work was undertaken with a view to filling in the gaps in the existing literature on the recrystallization behaviour of B-doped polycrystalline Ni3Al alloys.
free surface of the flat samples were chosen for this purpose. The annealed microstructures were characterized using a JEOL 2000 FX II transmission electron microscope (TEM) operated at 160 kV. The variation, during annealing, of order parameter (S) was monitored by X-ray diffraction technique. This parameter (S) was calculated from the integrated intensity ratios of the (100)/(200) and (110)/(220) pairs of reflections. The relationship used between S and the above intensity ratios is as follows [6]:
2. Experimental procedure The material used was the intermetallic alloy Ni76Al24 with 0.24 at.% boron. The arc melted alloy in the cold forged condition was obtained (courtesy Professor Dr. G. Gottstein of the Institut fu¨r Metallkunde und Metallphysik, RWTH Aachen, Germany). After homogenization at 1050°C for 25 min, the material yielded an average grain size of 30 mm. Homogenized samples of 4.5-mm thickness were cold rolled at room temperature to a reduction of 85%. Isochronal annealing of the cold rolled material was carried out from 100 to 1000°C at intervals of 100°C. The holding time was 1 h in each case. A few samples were also isothermally annealed at four different temperatures, namely 750, 800, 850 and 900°C for various lengths of time. Thin foils were prepared from the materials at different stages of annealing by mechanical polishing, followed by electropolishing in a jet electropolisher using a solution of 10% H2SO4 in methanol. Rolling plane sections at a depth of 1/4 the thickness from the outer
Fig. 2. Variation of the order parameter, S, with degree of cold rolling.
S. Ghosh Chowdhury et al. / Materials Science and Engineering A277 (2000) 1–10
Fig. 3. Variation of order parameter, S, during 1-h isochronal annealing of the 85% cold rolled material.
3
subscripts s and f refer to the superlattice and fundamental reflections, respectively. The rationale behind the choice of the (100)/(200) and (110)/(220) line pairs is that because the planes are parallel in any such pair, the value of the intensity ratio will not be affected by textural changes (although the absolute intensities of the lines will be obviously influenced by texture). For texture measurement, samples of 14×24 mm2 were cut out from the flat plates of different annealed materials. From each specimen, material was etched off from one side up to 1/4 of the thickness, and texture was measured on that surface only. In each case, a set of four pole figures, such as {111}, {200}, {220} and {311}, were measured and the orientation distribution function (ODF) was computed by means of series expansion [7] to an order of l= 22 and plotted in the f2 = constant sections.
3. Results and analysis
3.1. Structural changes
Fig. 4. Variation of the order parameter, S, with time at the isothermal annealing temperature of 850°C.
S2=
Is (xAl fAl +xNi fNi)2 (Lp)f (e − 2M)f 16 If ( fAl −fNi)2 (Lp)s (e − 2M)s
where, I is the measured integrated intensity; xAl and xNi are the atom fractions of Al and Ni, respectively, in the alloy; fAl and fNi are the atomic scattering factors for Al and Ni, respectively; Lp is the Lorentz-Polarization factor; and e − 2M is the Debye-Waller factor. The
3.1.1. Order parameter Fig. 1 shows the X-ray diffraction profile taken from the alloy, in the initial homogenized condition and after different degrees of cold rolling. The pattern for the homogenized alloy has been analyzed as that for an L12 structure, showing clear (100) and (110) superlattice peaks over and above the fundamental lines. With an increase in the degree of cold rolling, the (100) peak becomes weaker. The order parameter measured from (100)/(200) intensity ratio reduces to zero after 85% cold work. However, the intensity of the other superlattice peak (110) persists even after 85% cold rolling. In fact, the order parameter calculated from (110)/(220) intensity ratio gradually decreases and assumes a nearly constant value of 0.45 from 65% deformation onwards. The variations of the order parameter with amount of
Fig. 5. Variation of hardness with time during isothermal annealing at four different temperatures. Arrows indicate the start of recrystallization.
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The variations in the order parameter values during the course of 1-h isochronal annealing of the 85% reduced material are depicted in Fig. 3. It is observed that the intensity of the (100) peak remains zero even after isochronal annealing at 200°C. With the increase in temperature the value of the order parameter, S, measured from the (100)/(200) intensity ratio, increases rapidly until 400°C and then shows a gradual increase and attains the maximum value of 0.8 at 1000°C (Fig. 3). By contrast, the value of S from the (110)/(220) line pair shows a gradual increase with temperature to the initial value (before cold rolling) of 0.8. The above isochronal annealing curves clearly indicate a major change in order parameter between 200 and 400°C in the 85% cold rolled sample. The variation of S during isothermal annealing is shown typically in Fig. 4 for the sample annealed at 850°C. The variation of hardness during isothermal annealing at four different temperatures as depicted in Fig. 5 clearly shows that nearly 66% of the hardness recovers prior to the beginning of the recrystallization process.
3.1.2. Microstructure The cold rolled (85%) microstructure of this alloy was characterized by a high density of shear bands as clearly shown in typical optical micrographs (Fig. 6a). TEM micrographs have indicated the presence of a high density of dislocations, microbands as well as deformation twins (Fig. 6b). The SAD pattern from the twinned area in Fig. 6b is shown in Fig. 6c. Annealing at 500°C shows at many places groups of parallel stripes running nearly perpendicular to original grain boundaries (Fig. 7). Features similar to these stripes were observed in a Cu3Pt alloy during transformation from the disordered to the L12 structure [9]. Those stripes were identified by the authors as ordered domains of the L12 structure. It has already been seen from X-ray diffraction results (Fig. 3) that at this
Fig. 6. (a) Optical micrograph showing shear bands. (b) TEM micrograph showing high densities of dislocations, microbands and deformation twins. (c) SAD pattern of twinned area in (b).
rolling reduction have been plotted in Fig. 2. The above behaviour has been found to be consistent with a structural transformation from L12 to DO22 taking place in the material due to rolling deformation [8]. It might be mentioned that peak broadening due to shrinking crystallite size and increasing RMS strain during cold work is typical of essentially all metals. This fact may tend to ‘blur’ the superlattice peaks and hence, exaggerate the disordering. Perhaps the lower modulus value along B 100 \ contributes to the distinction between the two families of planes.
Fig. 7. TEM micrograph of stripe like features emanating from the grain boundaries after annealing at 500°C (Moire´ fringes like features near a grain boundary).
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Fig. 9. TEM micrograph of a group of grains after annealing at 700°C.
Fig. 10. TEM micrograph of a partly recrystallized area after annealing at 800°C.
Fig. 8. (a) TEM micrograph of twin like features after annealing at 500°C, (b) SAD pattern of the same area and (c) key to the SAD pattern.
temperature, the (100) superlattice peak which disappeared after heavy rolling deformation, reappears and the structure becomes L12 again. Presumably, annealing leads to a retransformation from the DO22 to the L12 structure. A distinct Moire´ fringe like feature sometimes observed at grain boundary regions (shown with an arrow in Fig. 7) can be speculated to indicate the possible coexistence of the above two structures. Twin like features are also sometimes observed in regions which have recovered sufficiently. This is shown typically in Fig. 8a. The SAD taken from this area (Fig. 8b) shows the [011] pattern of single crystal of the L12
Fig. 11. f2 sections of the ODF of 85% cold rolled boron-doped Ni76Al24.
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Fig. 12. f2 sections of the ODF after annealing at 850°C for 17 h.
tallization aids in the formation and growth of the Goss oriented regions. The ODF of the alloy, annealed at 1025°C for 1 h, again shows a Goss component along with rotated cube orientations (Fig. 13c). The texture components reported by different investigators (including the present work) during recrystallization of heavily rolled Ni76Al24(B) have been tabulated in Table 1. The variation of volume fraction recrystallized with time at various annealing temperatures is shown in Fig. 14 for the 85% cold rolled material [10]. As shown in Fig. 14 the alloy had yet to start recrystallizing after 1 h of annealing; only partial recrystallization set in after annealing for 5 and 11 h, whereas 24-h annealing produced slight grain growth beyond completion of recrystallization. In order to understand the effects of the structural change and the change in order during annealing on texture, measurements were made on samples isothermally annealed at 850°C for 1, 5, 11 and 24 h. The
structure but does not show distinct twin spots. The key to the diffraction pattern is given in Fig. 8c. Annealing at 700°C leads to copious nucleation of large number of fine grains (Fig. 9). Fig. 10 again shows a typical partly recrystallized area showing recrystallization twins inside. The SAD from the central grain shows the [011] zone axes of the L12 structure.
3.2. De6elopment of texture Texture measurements were undertaken on the partly and fully recrystallized samples which were initially cold rolled by 85%. As reported earlier [3], the texture of the 85% cold rolled material (Fig. 11) was reasonably sharp (maximum intensity 5 R). The texture after primary recrystallization has been found to be invariably weaker (maximum intensity 1.5 R). Fig. 12 shows the ODF of a fully recrystallized (metallographically) sample, annealed at 850°C for 1 h. A comparison between the above two figures clearly indicates that the texture severity drastically decreases upon recrystallization and that the recrystallization texture of this material is very weak. Even then some distinct features are discernible in the ODF of the recrystallized sample, such as the presence of a rotated cube {025}B 100\ along with a weak Goss {011} B100 \ component. The ODF of the sample recrystallized at 950°C for 1 h, however, shows only the weak rotated cube orientation, but no Goss orientation (Fig. 13a). The Goss orientation develops a bit later in this material, after 48 h of annealing (Fig. 13b) and there is a considerable amount of scatter around this orientation. The intensity of this orientation in this sample is also markedly higher than its intensities in the ODFs of the earlier samples. This clearly indicates that grain growth after primary recrys-
Fig. 13. (a) f2 =0° sections of the ODF after annealing at 950°C for 1 h. (b) f2 =0° sections of the ODF after annealing at 950°C for 48 h. (c) f2 =0° sections of the ODF after annealing at 1025°C for 1 h.
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Table 1 Effect of temperature of annealing on the development of recrystallization texture components in Ni76Al24(B) Material
Temperature (°C)
Time
Component
Intensity recrystallization/roll
Reference
90% Cold rolled
650 700 750 710 750 850 950 950 1025
1 1 1 90 30 17 1 48 1
{310}B135\ {211}B135\ {012}B121\, as above {013}B100\, {025}B100\, {025}B100\ {011}B100\ {025}B100\,
3.3/16.2 3.8/16.2 2.0/6.3 3.4/6.3 2.0/6.8 1.6/5.0 1.5/5.0 1.8/5.0 1.6/5.0
[1] [1] [4] [4] [4] Present Present Present Present
92% 90% 92% 85%
Cold Cold Cold Cold
rolled rolled rolled rolled
h h h min min h h h h
{112}B131\ {102}B201\, {112}B294\ {011}B100\
{011}B100\
work work work work
Fig. 14. Variation of the volume fraction recrystallized with time at various annealing temperatures.
f2 = 0° sections of the ODFs of these four samples are presented in Fig. 15a – d. All these ODFs are found to be essentially similar in the sense that all are very weak, much weaker than the ODF of the 85% cold rolled material (Fig. 11). The reproducible textural components after 1 h of annealing are {025} B 100 \ and {102}B 201\. After 5 h, a weak (near) Goss component {011}B 311\ is also observed in addition to the above components. After 11 h of annealing, all the previous components are obtained, but here the (near) Goss component {011}B 311 \ is found to be the most intense and has the exact Goss component also in its spread. The {025}B100 \ component is non-existent after 24-h annealing, indicating that the (near) Goss component of texture survives best on continued annealing. The development of recrystallization texture in boron-doped Ni76Al24 alloy is very different from that of pure Ni which is known to exhibit a strong cube {100}B 001\ texture under identical conditions. Both the materials have nearly the same stacking fault energy (SFE) as mentioned in Refs. [1] [2]. Irrespective of that, the deformation as well as recrystallization textures have been found to be widely different in the two cases. Presumably, factors such as ordering of Ni76Al24(B) and
the structural changes it undergoes during deformation (L12 DO22) as well as during annealing (DO22 L12) far outweigh the effect of SFE in determining the final textures. The 1-h annealed sample represents only a recovery stage. However, the S parameter was found to
Fig. 15. f2 =0° sections of the ODFs after annealing at 850°C: (a) 1 h, (b) 5 h, (c) 11 h and (d) 24 h.
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Fig. 16. XRD line profiles for samples cold rolled 85% and then isothermally annealed at 850°C for 5 and 30 min.
Fig. 17. DSC thermogram of 85% cold rolled boron-doped Ni76Al24.
increase from zero (after 85% cold rolling) to a value of 0.64 at this stage (Fig. 4). The XRD line profile for the samples annealed at 850°C for 5 and 30 min as well as for the 85% cold rolled sample have been plotted in Fig. 16. Clearly, a structural reordering from the DO22 to L12 structure has taken place during the above
annealing stages, which do not however, show any perceptible recrystallization. Thus, in a way, the formation of the weak annealing texture from a strong deformation texture can be termed as a kind of ‘transformation’ texture produced by the structural change from DO22 back to L12 during annealing.
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4. Discussion On the basis of what has been stated above, the annealing process in this alloy appears to involve a combination of three different phenomena: recovery, reordering and recrystallization. Several aspects need to be considered in explaining the results of the recrystallization texture measurements. The first pertains to the fact that a very weak texture is produced from the sharp cold rolling texture during the ‘recovery’ stage, which does not involve any reorientation of grains as a whole, but during which substantial atomic shuffling takes place producing the parent L12 structure from the DO22, which was the end structure after cold rolling. It is known that the DO22 structure can be derived from the L12 by the introduction of a 1/2[110] APB on every {001} plane. Similarly, the introduction of an APB on every (001) plane in the DO22 structure again transforms it back into the L12 structure [11]. The drastic weakening of the texture intensity after 1-h annealing at 850°C [ f(g)= 1.7] from that of the 85% cold rolled alloy [ f(g)= 5.0] can not be related to any reorientation of grains (since no recrystallization takes place at this stage), but, on the other hand, could be related to the DO22 L12 reverse transformation. It may also be recalled that annealing has been found to be associated with the appearance of a large density of twin like features. Tanner and Ashby [12,13] made a theoretical analysis of the relief of strains associated with certain phase transformations in intermetallic alloys, such as Ni3V which has a DO22 structure at room temperature. According to them, when ordering in a binary alloy involves a change in crystal symmetry the reaction is invariably accompanied by an accumulation of substantial internal strains. Under certain conditions, these strains can be relieved through twinning of the product phase. Thus, the twin like features obtained in the product phase after annealing (long before recrystallization takes place) in the present boron-doped Ni76Al24 alloy (Fig. 8a) can be assumed to be due to such an effect. It has been shown [14] that during annealing, the heat evolution before the onset of recrystallization is quite substantial (11.5 kJ/mol) (peak A in Fig. 17). In comparison to that, the energy expended during recovery for Ni varies from 3.7 to 5.4 kJ/mol [15]. Therefore, a major part of the stored energy release in the present material can be attributed to the DO22 L12 reverse transformation. This clearly indicates that significant relief of strain is associated with this transformation, and it is no wonder, therefore, that a high density of twin like features accompanies this transformation. The incidence of twinning will further lead to fragmentation of the grains, resulting in a weakening of the texture intensity after 1 h an.
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nealing at 850°C from that of the 85% cold rolled alloy. The weak texture obtained after annealing for 1 h at 850°C continues as it is during progressive annealing up to 5, 11, 17 and 24 h, i.e. during the stages of partial recrystallization to full recrystallization, continuing even up to the grain growth stage. Very weak recrystallization texture in Ni3Al(B) has also been reported by Ball and Gottstein [2]. The almost random recrystallization texture in this alloy has been attributed by them to the high nucleation rate and the low grain boundary mobility, which produces a lack of growth selection during recrystallization. Using this argument it is difficult to reason out why the recrystallization texture should be so very weak compared to the cold rolling texture and at the same time why it should not bear any resemblance at all with the essential features of the cold rolling texture. On the other hand, the results of the present investigation suggest that when recrystallization starts during annealing, the sharp cold rolling texture has already been replaced by a weak ‘transformation’ texture. Subsequent recrystallization taking place in such a weakly textured matrix is expected to give rise to only a weak recrystallization texture. Of all the orientations observed in the ODFs of the alloy during partly or fully recrystallized conditions, the Goss component {011}B 100\ appears to be the most prominent. It is known that this component usually originates in deformation/shear bands. In the present alloy, formation of shear bands started from 65% cold rolling onwards and the density of such bands increased significantly after 85% cold work [3].
5. Conclusions The DO22 structure (the end product after 85% cold rolling) reverts back to the initial L12 during annealing. The reordering process is accompanied by the formation of twins which might accommodate the transformation strain. During annealing the texture changes drastically even before the onset of recrystallization and becomes very weak, almost random. This can be correlated with the DO22 L12 transformation. The formation of twins during the transformation leads to fragmentation of the matrix resulting in a weak texture. During recrystallization a weak recrystallization texture is produced from the already recovered and reordered material which itself has a weak texture. The prominent recrystallization texture components are {025}B100\ and {011}B 100\ . The Goss component becomes prominent in the post-recrystallization stage.
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Acknowledgements The authors are indebted to Professor G. Gottstein, Director, Institute fu¨r Metallkunde und Metallphysik, RWTH Aachen, Germany for the provision of laboratory facilities for conducting the texture work. Acknowledgements are also due to V. Kumar and Mr. Umashankar for their help during the experimental work.
References [1] G. Gottstein, P. Nagpal, W. Kim, Mater. Sci. Eng. A108 (1989) 165. [2] J. Ball, G. Gottstein, Intermetallics 1 (1993) 171. [3] S. Ghosh Chowdhury, R.K. Ray, A.K. Jena, Mater. Sci. Eng. A A246 (1998) 289.
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[4] J. Ball, G. Gottstein, Intermetallics 1 (1993) 191. [5] J. Ball, G. Gottstein, Intermetallics 2 (1994) 205. [6] B.D. Cullity, Elements of X-Ray Diffraction, Narosa, London, 1993. [7] H.J. Bunge, Texture Analysis in Materials Science, Butterworth, London, 1982. [8] S. Ghosh Chowdhury, R.K. Ray, A.K. Jena, Scripta Metall. Mater. 32 (1995) 1501. [9] K. Mitsui, Y. Mishima, T. Suzuki, Phil. Mag. A53 (1986) 357. [10] S. Ghosh Chowdhury, A.K. Jena, R.K. Ray, Metall. Mater. Trans. A 29A (1998) 2893. [11] M. Yamaguchi, Y. Umakoshi, in: R.W. Cahn, P. Hassen, E.J. Kramer (Eds.), Plastic Deformation and Fracture of Solids, VH, Weinheim, 1993, p. 251. [12] L.E. Tanner, Phys. Stat. Sol.(a) 30 (1968) 685. [13] L.E. Tanner, M.F. Ashby, Phys. Stat. Sol. 33 (1969) 59. [14] S. Ghosh Chowdhury, Ph.D Thesis, I.I.T. Kanpur, India, 1995. [15] L.M. Clarebrough, M.E. Hargrieves, G.W. West, Proc. R. Soc. A232 (1955) 252.