- Email: [email protected]
Microstructural evolution of commercial pure iron during directional annealing
Materials Science and Engineering A 422 (2006) 241–251
Microstructural evolution of commercial pure iron during directional annealing Z.W. Zhang a,∗ , G.L. Chen a,b , G. Chen a a
Joint Laboratory of Nanostructured Materials and Technology, Nanjing University of Science and Technology, Nanjing 210094, PR China b State key lab for advanced metals and materials, USTB, Beijing 100083, PR China Received 26 January 2006; accepted 3 February 2006
Abstract Microstructural evolution during the directional annealing of the cold-rolled pure iron has been investigated. Columnar grains with the largest aspect ratio of 23.6 were produced in the specimen cold-rolled for 70% of thickness reduction at the hot zone temperature (HZT) of 850 ◦ C and a temperature gradient (TG) of 200 ◦ C/cm with an optimum drawing velocity of 5 m/s. It was found that the columnar grain structure was developed due to abnormal grain growth. There is a critical coarsening temperature, below which the secondary recrystallization does not occur during the annealing; consequently no columnar grain structure can be produced. No obvious evidence shows that the grain size after primary recrystallization has an influence on the microstructural evolution. Due to the strong crystallographic texture of the small grains ahead of the hot zone, the columnar grains have small aspect ratios. Most of the boundaries of the columnar grain are high angle coincidence site lattice (CSL) grain boundaries. The grains with twin or low angle grain boundaries are apt to form island grains in the columnar structure. © 2006 Elsevier B.V. All rights reserved. Keywords: Directional annealing; Columnar-grained structure; Commercial pure iron; Texture
1. Introduction Columnar-grained microstructures, with grain boundaries aligned parallel to the applied tensile stress direction, can enhance creep properties [1–3], improve fatigue resistance [4] and impart crack-stopping behaviors [5]. These structures can often be produced by directional solidification. It may be desirable, however, to use material processed in the solid state. In those cases, directional recrystallization is the possible processing route. Directional recrystallization (zone annealing) is the annealing process whereby the sample is traversed through a furnace (hot zone) containing a steep temperature gradient (with the temperature varying along a direction parallel to the traverse direction) [4,6]. The microstructural evolution during directional recrystallization (DR) depends sensitively on materials and the preparation parameters. In many earlier studies, the DR technique was mainly used on processing superalloys, especially the oxide
∗
Corresponding author. Tel.: +86 025 84315797; fax: +86 025 84315797. E-mail address: [email protected] (Z.W. Zhang).
0921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2006.02.001
dispersion strengthened (ODS) nickel and iron-based superalloys for applications in the aerospace industry [7–13]. For ODS superalloys, the prerequisite for the successful DR is a fine (∼0.2 m) equiaxed matrix grain size [7,8]. In some superalloys, simple static annealing will result in large elongated grains because the initial powder consolidation and rolling or extrusion lead to the alignment of particles which act to pin grain boundaries anisotropically [9–11]. While the principle of such zone annealing has been known for a long time, the process parameters of temperature gradient (TG), hot zone temperature (HZT) and drawing velocity, together with material characteristics, must be closely controlled to produce the desired structure [9]. Humphreys [8] have shown that, in directionally recrystallized MA-6000, drawing velocity can influence the microstructure and the crystallographic texture of MA-6000 extrusions. This can be rationalized using the difference in boundary mobility due to the solute pinning. Recently, Li et al. [15] and Baker and Li [14] researched systematically the effect of processing parameters of the directional annealing on the directional recrysatallization of poly crystalline nickel and cold-rolled copper single crystal. They observed that the processing parameters also influence the number of annealing
242
Z.W. Zhang et al. / Materials Science and Engineering A 422 (2006) 241–251
twins per unit of area and texture evolution, in addition to the microstructure. This work is a part of the study aiming to develop a fundamental understanding of how the initial microstructural parameters and the processing variables determine the microstructural evolution of commercial pure iron during directional annealing. 2. Experimental work The material used in the present work was in the form of bars with an average grain size of 50 m, as determined by the linear intercept method. The microstructure of the as-received materials is shown in Fig. 1. The composition of the commercial pure iron is shown in Table 1. Specimens were cold-rolled to 70% and 85% thickness reduction, with thickness reductions limited to 0.5 mm per pass (<2% of thickness reduction) during the last three passes. The specimens were water-cooled after each pass to prevent heating. The strips after cold rolling for 70% and 85% thickness reduction were 8 and 4.5 mm thick, respectively. And then the strips were cut into bars, which have square crosssections of 8 × 8 mm2 and 4.5 × 4.5 mm2 and are more than 150 mm long. The rolled bars were polished mechanically and then cleaned using acetone. The recrystallization temperatures of the rolled specimens were measured by differential thermal
Fig. 1. Microstructure of as-received commercial pure iron. Table 1 Composition of commercial pure iron in weight percent Element
wt.%
C Si Mn P S Al Ni Cr Cu Ti Fe
0.003 0.025 0.150 0.010 0.010 0.040 / / / / Bal.
Fig. 2. Schematic illustration of the directional annealing furnace. The induction coil is held stationary and the specimen moved from bottom to top or inversely.
analysis (DTA) with a Shimadzu DTA-50 thermal analysis system in nitrogen gas. The specimens were heated from ambient to 1000 ◦ C at a heating rate of 20 ◦ C/min. The measured recrystallization temperatures were 662 and 640 ◦ C for the specimen cold-rolled for 70% and 85% of reduction, respectively. The experimental set-up for the directional annealing is shown in Fig. 2. A 5 mm high single circle induction coil and a Ga-In liquid alloy cooling pool are used. The specimen is held on a transition bar (drawing bar), which is moved through a collet by a stepper motor at speeds ranging from 0.1 m/s to 9000 m/s. The distance between the surface of the Ga–In liquid alloy and the induction coil can be adjusted to change the temperature gradient (TG) ahead of the hot zone. In the present study, the TG was 200 ◦ C/cm and the hot zone temperatures (HZTs) of 700, 750 and 850 ◦ C were used for directional annealing. In several studies of the directional annealing of the nickelbased superalloy INCONEL MA6000 [5], cold-rolled copper single crystals [14] and polycrystalline nickel [15], the specimen was run once for several centimeters at a given temperature but at different moving speeds. In order to distinguish these regions for different speeds, the interval regions of short lengths were run at a very slow rate and a low temperature. As a result, the obtained strip possessed several experimental “runs” with small unrecrystallized intervals. The same technique was employed here. However, in our experimental process, the interval regions were run at a very high rate and a low temperature. With this modification, the same result was obtained; and meanwhile the experimental time was reduced substantially. After directional annealing, the specimens were mechanically polished and etched in 4% nitric acid in ethanol for optical metallography. Grain sizes were measured using the linear intercept method. For specimens with columnar grain structures, grain size measurements were made both parallel and perpendicular to the growth direction. The preferred crystallographic orientations of the as-rolled and the directional-annealed materials were determined by electron backscattered diffraction (EBSD) and analysis of X-ray
Z.W. Zhang et al. / Materials Science and Engineering A 422 (2006) 241–251
reflection pole figures. A Siemens 5000 X-ray diffractometer was employed to measure the textures by recording the 110, 200 and 211 pole figures from longitudinal sections. The Xradiation used was CuK␣ with a wavelength of 0.15406 nm. The EBSD measurements were carried out using a Leo1450 scanning electron microscopy (SEM) equipped with HKL Channel 4.2 software. And TexTools v.3.2, commercial software (ResMat Corp.) was also used for texture analysis. The value of coincidence site lattice (CSL) grain boundary was calculated using Frank-Ranganthan method [16]. 3. Experimental results 3.1. Deformation degree Figs. 3 and 4 show the microstructures of the specimens coldrolled for 70% and 85% of thickness reduction, respectively. The specimens were directionally annealed at 850 ◦ C with various drawing velocities. Obviously, the recrystallized microstructures show different features under different degree of deformation and drawing velocities. At the drawing velocity of 1 m/s, the microstructure consists of large grains, between which there are
243
some small grains (Fig. 3 (a)). Fig. 3(b) shows a columnar grain structure, which was developed at a drawing velocity of 3 m/s. The longitudinal grain size is approximately 6.7 mm and the transverse grain size is 600 m. Increasing the drawing velocity to 5 m/s, the longitudinal and transverse grain sizes are approximately 11.8 mm and 500 m, respectively. The corresponding micrograph is shown in Fig. 3(c). It is worth noting that the boundaries between the large grains grown at the drawing velocities of 3 and 5 m/s are almost parallel to the movement direction of the specimen. As the drawing velocity was increased further to 25 m/s, a small-sized equiaxed grain structure formed with an average grain size of approximately 50 m (Fig. 3(d)). For the degree of deformation of 85% reduction, directional recrystallization occurred at drawing velocity of 3–8 m/s, as shown in Fig. 4. However, the longitudinal grain sizes are smaller and the transverse grain sizes are almost the same comparing with the specimens, which are cold-rolled for reduction of 70%. In addition, we noticed that, for the specimens cold-rolled for 85% reduction, the range of drawing velocity for the occurrence of the directional recrystallization is broader, from 3 to 8 m/s. Fig. 5 shows the profile of the relationship between the grain sizes and the drawing velocities. In order to delineate the grain structure
Fig. 3. Microstructures of commercial pure iron cold-rolled for 70% of thickness reduction and directionally annealed at 850 ◦ C with a temperature gradient of 200 ◦ C/cm and various drawing velocity of (a) 1 m/s, (b) 3 m/s, (c) 5 m/s and (d) 25 m/s. The arrows refer to the direction of columnar grain growth.
244
Z.W. Zhang et al. / Materials Science and Engineering A 422 (2006) 241–251
Fig. 4. Microstructures of commercial pure iron cold-rolled for 85% of thickness reduction and directionally annealed at 850 ◦ C with a temperature gradient of 200 ◦ C/cm and various drawing velocity of (a) 3 m/s, (b) 5 m/s, (c) 8 m/s and (d) 25 m/s. The arrows refer to the direction of columnar grain growth.
more precisely, the grain sizes we used in Fig. 5 were the sizes perpendicular and parallel to the columnar growth direction, separately, instead of the average grain sizes. 3.2. Hot zone temperature (HZT)
Fig. 5. Grain size as a function of drawing velocity for the degree of deformation 70% and 85% reduction at the hot zone temperature of 850 ◦ C and the temperature gradient of 200 ◦ C/cm. When a columnar grain structure is present, the grain size is plotted not as the average grain size but as the grain size perpendicular and parallel to the columnar growth direction.
Fig. 6 shows the grain sizes as a function of drawing velocities for the two HZTs of 700 and 750 ◦ C. Fig. 7 shows the microstructures of the specimens cold-rolled for 85% of thickness reduction. The specimens were directionally annealed at 700 ◦ C with the drawing velocities of 1 m/s and 8 m/s. Smallsized equiaxed grain structures are observed at all the drawing velocities from 1 to 8 m/s. At the drawing velocity of 1 m/s, the average grain size is about 50 m (Fig. 7 (a)). Increasing the drawing velocity, the average grain size decreases slightly. As the drawing velocity was increased further to 8 m/s, a partially recrystallized structure was produced with an average recrystallised grain size of approximately 20 m (Fig. 7(b)). Fig. 8 shows the microstructures of the specimens cold-rolled for reduction of 70%. The specimens directionally annealed at 750 ◦ C with the drawing velocities of 1, 3 and 5 m/s. It can be seen clearly from Fig. 8 (b) that the columnar-like grain struc-
Z.W. Zhang et al. / Materials Science and Engineering A 422 (2006) 241–251
245
ture can be developed only at the drawing velocity of 3 m/s. However, The longitudinal grain size of the elongated grains is much smaller (1.2 mm) than that directionally annealed at 850 ◦ C, although the transverse grain size is almost the same (330 m). At both drawing velocities of 1 and 5 m/s, the grain structures are equiaxed with the average grain size of about 100 m, see Fig. 8 (a) and (c). 3.3. Grain size after primary recrystallization
Fig. 6. Grain size as a function of drawing velocity for the hot zone temperature (HZT) of 700 ◦ C and 750 ◦ C at the temperature gradient of 200 ◦ C/cm. When a columnar grain structure is present, the grain size is plotted not as the average grain size but as the grain size perpendicular and parallel to the columnar growth direction.
To test the effect of grain sizes after primary recrystallization on the columnar grain growth, pure iron specimens cold-rolled for 70% of thickness reduction were isothermally annealed at 650, 700 and 750 ◦ C for 30 min. After isothermally annealed at 650 ◦ C, the specimen had a partially recrystallized structure. The specimens isothermally annealed at 700 ◦ C and 750 ◦ C had completely recrystallized structures with grain size of 50 and 100 m, respectively. These specimens were then directionally annealed at 850 ◦ C and a TG of 200 ◦ C/cm with the drawing velocity of 3 m/s. The resulting microstructures are shown in Fig. 9. The resulting microstructures are very similar to those shown in Fig. 3, which were directionally annealed after cold rolling without isothermal annealing. No obvious evidence indicated that the grain size after primarily recrystallization influences the microstructural evolution. This result supports the suggestion that the columnar grains in the cold-rolled commercial pure iron were produced by the secondary recrystallization. This, perhaps, is not surprising. Under the temperature gradient of 200 ◦ C/cm ahead of the hot zone, at 5 mm ahead of the hot zone the temperature is approximately 700 ◦ C. This is higher than the temperature (662 ◦ C), which allows the commercial pure iron to recrystallize. 3.4. Crystallographic texture
Fig. 7. Microstructures of commercial pure iron cold-rolled for 85% of thickness reduction and directionally annealed at 700 ◦ C and a temperature gradient of 200 ◦ C/cm with drawing velocities of (a) 1 m/s and (b) 8 m/s The direction of annealing was from left to right.
The preferred crystallographic orientations of as cold-rolled texture were determined using data from X-ray reflection pole figures. The orientations of the primary recrystallised small grains ahead of the hot zone and those of the large columnar grains were determined from EBSD attached to a SEM. The textures induced by various processing parameters are summarized in Table 2. Under the cold-rolled condition, the materials had an elongated grain structures. The deformation textures, centered on <1 1 0> fiber with {1 1 1}<1 1 0> and {0 0 1}<1 1 0> components, were developed after reduction of 70% and those, centered on <1 1 0> fiber with {1 1 2}<1 1 0> and {1 1 1}<1 1 0> components, after larger thickness reduction of 85%. When the specimen cold-rolled for 85% of reduction was directionally annealed at 700 ◦ C with the drawing velocity of 1 m/s, recrystallization occurred, resulting in a microstructure with small-sized equiaxed grain, see Fig. 7 (a). Under this condition, this primary recrystallized microstructure had textures, which were different in type from the corresponding cold-rolled textures. There was a tendency for {1 1 2}<1 1 0> component to be depleted. Most primary recrystallized textures were dominated by the near {1 1 1}<1 1 0> component.
246
Z.W. Zhang et al. / Materials Science and Engineering A 422 (2006) 241–251
Fig. 8. Microstructures of commercial pure iron cold-rolled for 70% of thickness reduction and directionally annealed at 750 ◦ C with a temperature gradient of 200 ◦ C/cm and various drawing velocity of (a) 1 m/s, (b) 3 m/s and (c) 5 m/s. The direction of annealing was from left to right.
For the specimens cold-rolled for 85% of reduction and directionally annealed at 850 ◦ C with the drawing velocity of 5 m/s, the recrystallized textures of the small grains ahead of the hot zone are distinctly different from the deformation textures. Most recrystallized textures of the small grains were dominated by the near {1 1 0}<1 1 2> component. Also very weak {1 1 0}<1 1 0>
and {1 1 1}<1 2 1> components are shown in these materials. However, for the specimen cold-rolled for 70% of thickness reduction and directionally annealed with the same processing parameters as those cold-rolled for 85% of thickness reduction, the small grains ahead of the hot zone have a random texture with very weak {1 1 1} <1 2 1> and {1 1 0}<1 1 2> components.
Table 2 Crystallographic textures of commercial pure iron under various processing conditions Degree of deformation (%)
Process variables
Texture
Grain structure
70
As cold-rolled
Elongated grains
85
Rolled and directionally annealed at 850 ◦ C As cold-rolled Rolled and directionally annealed at 700 ◦ C
Stronger components {1 1 1}<1 1 0>, {0 0 1}<1 1 0> Weaker components {1 1 0}<1 1 0>and {1 1 1}<2 1 1> Random texture with very weak {1 1 1} <1 2 1> and {1 1 0}<1 1 2> components Stronger components {1 1 2}<1 1 0> and {1 1 1}<1 1 0>, Weaker components {1 1 0}<1 1 0>, {1 1 1}<2 1 1> <1 1 0>fiber with stronger {1 1 1}<1 1 0> and weaker {1 1 1}<1 2 1>components
Rolled and directionally annealed at 850 ◦ C
Stronger {1 1 0}<1 1 2>, weaker {1 1 0}<1 1 0>and {1 1 1}<1 2 1> components
Columnar grain with high aspect ratio Elongated grains Equiaxed fine grains (Primary recrystallization) Columnar grain with low aspect ratio
For the cold-rolled specimen, texture refers to deformation texture; for those cold-rolled and directionally annealed, textures refer to primarily recrystallized texture of the small grains ahead of the hot zone; Grain structures refer to those after directional grain growth except the as cold-rolled specimens.
Z.W. Zhang et al. / Materials Science and Engineering A 422 (2006) 241–251
247
Fig. 9. Microstructures of commercial pure iron cold-rolled for 70% of reduction and then directionally annealed at 850 ◦ C and a temperature gradient of 200 ◦ C/cm with the drawing velocity of 3 m/s. Before directionally annealed, the cold-rolled specimens were isothermally annealed for 30 min at (a) 650 ◦ C, (b)700 ◦ C and (c) 750 ◦ C. The arrow refers to the direction of columnar grain growth.
We noticed that the specimens cold-rolled for 85% of reduction had the same texture with {1 1 1}<1 1 0> component both when directionally annealed at 850 ◦ C with the drawing velocity of 1000 m/s and directionally annealed at 700 ◦ C with the drawing velocity of 1 m/s, although the former had a smaller grain size of approximately 35 m. Fig. 10 shows the ODF section at Euler angle ϕ2 = 45◦ of the specimens cold-rolled for 85% of reduction and directionally annealed
with the two different processing parameters. This indicates that during the primary recrystallization, the recrystallized textures formed by directional annealing are similar to those by isothermal annealing. Moreover, from comparing the textures of the small equiaxed grains ahead of the hot zone induced by different processing parameters, it is apparent that the recrystallization textures show some dependence on the degree of deformation and heat treatment.
Fig. 10. Section through the ODF at Euler angle 2 = 45◦ of the specimen cold-rolled for 85% of reduction and directionally annealed (a) at 700 ◦ C and a temperature gradient of 200 ◦ C/cm with the drawing velocity of 1 m/s and (b) at 850 ◦ C and the temperature gradient of 200 ◦ C/cm with the drawing velocity of 1000 m/s. The contour levels shown are multiples of random density. The dominant texture components occur close to a line with = 54.7◦ , which corresponds to {1 1 1} lying parallel to the sheet plane. The texture is quite typical of recrystallized ferritic steel plate.
248
Z.W. Zhang et al. / Materials Science and Engineering A 422 (2006) 241–251
Fig. 11 shows the microstructures of the large columnar grains in the specimens cold-rolled for 85% and 70% of thickness reduction. The specimens were directionally annealed at 850 ◦ C with the drawing velocity of 5 m/s. The corresponding orientations of the columnar grains are listed in Table 3. It can be seen from Fig. 11 and Table 3 that the columnar grains after directional annealing have random orientations for both specimens cold-rolled for 85% and 70% of thickness reduction. However, the strong recrystallized texture {1 1 0}<1 1 2> of small grains in the specimen cold-rolled for 85% of thickness reduction produces some columnar grains with the orientation of {1 1 0}<1 1 2> (Grain 3 in Fig. 11 (a)), the recrystallized texture, although of course the statistics of these grains are limited. Under this condition, the columnar grains with smaller aspect ratio were produced. The relationships of the orientation between different columnar grains shown in Fig. 11 are summarized in Table 4. It can be seen that most of the misorientations between the columnar grains are concentrated on high angle coincidence site lattice (CSL) boundaries. Moreover, almost all the island grains (Grain4, 6, 8,12 in Fig. 11 (a) and grain 4, 7 in Fig. 11 (b)) have the twin grain boundaries (60◦ /<1 1 1>)or low angle grain boundaries with respect to the matrix columnar grains. 4. Discussion 4.1. Grain boundary migration velocity
Fig. 11. Microstructures ahead of the hot zone of the specimens cold-rolled for (a) 85% of thickness reduction and (b) 70% of thickness reduction and then directionally annealed at 850 ◦ C with a temperature gradient of 200 ◦ C/cm and the drawing velocity of 5 m/s. The arrows refer to the direction of columnar grain growth.
Microstructural evolution typically involves the movement of grain boundaries within the material. The movement velocity of a grain boundary is given by the well-known rate equation [17–19]. ν = MG where G is the driving force, usually expressed as the weighted mean curvature of the boundary during abnormal grain growth
Table 3 The orientation of the columnar grains in the specimens cold-rolled for thickness reduction of 85% and 70% and directionally annealed at 850 ◦ C and a temperature gradient of 200 ◦ C/cm with the drawing velocity of 5 m/s 85% of thickness reduction Grain no.
1 2 3 4 5 6 7 8 9 10 11 12
70% of thickness reduction
Test Euler angle
(◦ )
(hkl)[uvw]
∅1
∅2
0.4 222.1 325.1 216.9 320.0 217.7 318.8 315.2 58.3 0.4 88.4 0.9
8.3 19.2 45.3 34.2 27.2 33.1 28.5 28.6 37.1 1.2 45.5 27.5
42.6 60.5 91.3 22.9 73.4 21.8 72.4 77.2 28.8 179.2 68.8 1.2
The corresponding microstructures were shown in Fig. 11.
(119)[1-10] (216)[14-1] (101)[1-2-1] (134)[-22-1] (317)[3-2-1] (134)[-22-1] (317)[3-2-1] (317)[3-2-1] (123)[1-53] (001)[-100] (212)[-3-24] (012)[100]
Grain No.
1 2 3 4 5 6 7 8
Test Euler angle (◦ )
(hkl)[uvw]
∅1
∅2
144.8 67.2 261.5 1.2 151.7 201.1 127.9 94.8
30.8 49.2 36.7 46.8 71.3 47.6 34.5 30.5
88.6 35.3 26.1 15.2 64.8 22.8 45.1 74.6
(102)[-231] (233)[0-11] (123)[14-3] (144)[4-10] (943)[-232] (255)[-10 7-3] (112)[-201] (316)[-9-35]
Z.W. Zhang et al. / Materials Science and Engineering A 422 (2006) 241–251
249
Table 4 Misorientation of the columnar grains in the specimens cold-rolled for thickness reduction of 85% and 70% and directionally annealed at 850 ◦ C and a temperature gradient of 200 ◦ C/cm with the drawing velocity of 5 m/s 85% of thickness reduction
70% of thickness reduction
Index1
Index2
Misorientation
GB type
Index1
Index2
Misorientation
GB type
1 2 2 3 4 5 6 7 7 9 10 11
2 3 4 4 5 6 7 8 9 10 11 12
39.1◦ /[0-1-1] 57.7◦ /[1-1-1] 44.8◦ /[-1-4-1] 44.5◦ /[4-4-1] 58.2◦ /[-11-1] 57.0◦ /[-11-1] 58.9◦ /[-11-1] 2.4◦ /[3-3-1] 55.7◦ /[2-1-2] 36.6◦ /[210] 49.8◦ /[5-21] 58.7◦ /[11-1]
9 3 59 229 3 3 3 LA 45 27 87 3
1 2 3 4 5 6 6 7
2 3 5 5 6 7 8 8
45.8◦ /[-121] 44.6◦ /[213] 38.7◦ /[40-1] 58.1◦ /[-1-11] 47.4◦ /[012] 59.2◦ /[-111] 50.6◦ /[-211] 18.1◦ /[30-1]
21 95 161 3 15 3 31 LA
The corresponding microstructures were shown in Fig. 11; Note: Index: the number of the columnar grains marked in the Fig. 11 (a) and (b); LA: low angle boundary.
and M is the grain boundary mobility. Thus, a boundary travels twice as fast when its mobility doubles. Abnormal grain growth is generally assumed to arise from a combination of the inhibition of normal grain growth and the occurrence of a small number of special grains with a low free energy, providing a net driving force for their growth [20]. The resulting texture is claimed to be dependent on the driving force, which selects the “special” grains. Abnormal grain growth implies that one type of grains has exceptional fast-moving boundaries. There are two possible causes: (1) A special driving force G for selected grains (thermodynamics-based selection) and/or (2) a difference in M for different grain boundaries (kinetics-based selection). The typical driving force can be calculated from [21,22] G = Gv + Gb + Gs where Gv is the driving force from energy stored in the deformed matrix in the form of dislocations, Gb is the driving force from decreasing of the grain boundary area and Gs is from the surface energy. As mentioned above, for abnormal grain growth in commercial pure iron during directional annealing, since the columnar grains were produced by secondary recrystallization, the stored energy Gv from cold rolling had been completely consumed by the primary recrystallization. In addition, since the size of the iron bars are 8 mm × 8 mm and 4.5 mm × 4.5 mm in cross-section, the driving force from surface energy can be neglected. Hence, it can be proposed that the only driving force for columnar grain growth is the reduction of grain boundary area. The driving force from reducing of the grain boundary area can be calculated from [23] 1 1 Gb = E − D1 D2 where E is the energy of the grain boundary between the two neighboring grains with radius D1 and D2 , respectively. It is generally accepted that the energy of the grain boundary is a function of misorientation. The famous relationship between them was first derived by Read and Shockley [24] for a simple
tilt boundary based on a dislocation model and is given in the following equation E = E0 θ(A − ln θ) where E is the grain boundary energy, θ is the misorientation angle, E0 and A are constants. The mobility of high angle grain boundaries is temperaturedependent and is often found to obey an Arrhenius equation M = M0 exp(−Q/kT ) where Q is the activation energy for grain boundary migration, k the Boltzmann constant, and M0 is the usual prefactor. As far as the directional annealing is concerned, a mobility gradient implies an effective temperature gradient [25]. In addition, even if under the same high temperature, the differences in mobility may be significant, especially in textured materials [26,27]. The high mobility of the high-angle grain boundaries will cause
Fig. 12. Grain size as a function of drawing velocity for the specimens cold-rolled for 70% of thickness reduction. The specimens were directionally annealed at the hot zone temperature of 750 ◦ C and 850 ◦ C and a temperature gradient of 200 ◦ C/cm with the drawing velocity of 5 m/s. When a columnar grain structure is present, the grain size is plotted not as the average grain size but as the grain size perpendicular and parallel to the columnar growth direction.
250
Z.W. Zhang et al. / Materials Science and Engineering A 422 (2006) 241–251
some “special” grains grow rapidly by a kinetics-based selection mechanism during directional annealing. 4.2. Effect of processing variables From the studies mentioned above, it is clearly that the microstructure evolution strongly depends on the processing parameters. Fig. 12 shows the effect of the HZT on the grain size and shape. At the same TG and the drawing velocities, the higher HZT (850 ◦ C) benefits the forming of columnar grain structures. This suggests that, for the secondary recrystallization, there is a critical temperature, which is also called temperature of grain coarsening. When the annealing temperature is below this temperature the abnormal grain growth cannot occur, so the large columnar grain structure cannot be developed, see Figs. 3 and 8. The aspect ratios of the columnar grains under various processing parameters were shown in Fig. 13. The degree of deformation is not the key factor in determining the columnar grain growth. The specimens cold-rolled for 85% of reduction have more stored energy than those cold-rolled for 70% of reduction, however, the aspect ratio of columnar grains in these materials is smaller. This, perhaps, is not surprising, because the columnar grains were produced not by primary recrystallization but by secondary recrystallization. During primary recrystallization the stored energy drives the nucleation and normal grain growth. After the primary recrystallization, the store energy is depleted. At least, the residual strain energy is not the key driving force for the directional grain growth. Drawing velocity has an important influence on the aspect ratio of columnar grains. There is an optimum range of drawing velocities for columnar grain growth. The interpretation of the influence of drawing velocities on the microstructural evolution during the directional annealing requires consideration of the two distinct velocities involved in columnar grain growth [19]: the drawing velocity and the grain boundary movement. Holm et al. [19] simulated the microstructural evolution during zone annealing with different hot zone moving velocity (expressed as
Fig. 13. Aspect ratio of grains as a function of drawing velocity under various processing parameters. When an equiaxed grain structure is present, the aspect ratio equals to 1.
drawing velocity in the present paper). At small drawing velocities, the grain boundaries in the region ahead of the hot zone move faster than the drawing velocity. Thus, the fastest evolution is due to normal grain growth in this region, with a little extra growth while the hot zone is present. This leads to a largesized, equiaxed microstructure throughout the sample, as shown in Fig. 3 (a). When the drawing velocity is equal to the velocity of the boundaries in the hot zone, columnar grain growth occurs. If the drawing velocity is faster than that of the boundaries, both the length and width of grains are approximately equal to their normal grain size because the hot zone sweeps through the specimen too quickly to influence the boundaries. 4.3. Effect of crystallographic texture The columnar grains, which contact their neighbors by highangle CSL boundaries, propagate continuously and preferentially along the direction of TG, as shown in Fig. 11 and Table 4. The selective high-angle CSL boundary migration can be explained by the higher impurity concentration in the general high-angle grain boundaries than in high-angle CSL boundaries, which prevented the movement of general high-angle grain boundaries. It has been known [28] that the effects of impurities on boundary mobility are dependent on the crystallography of the boundary. The special boundaries, i.e. those which are close to a coincidence relationship, are less susceptible to effects of impurities than random boundaries. According to the findings reported by Messina et al. [29] and Park et al. [30], 3 and5 boundaries are expected to have low energy and low mobility. It is easy for these kinds of boundaries to form island grains, growing out into neighboring columnar grains. The similar result was found in this study. As noted above, directional secondary recrystallization of heavily cold-rolled commercial pure iron produced columnar grains with smaller aspect ratio than those in specimens coldrolled for 70% of thickness reduction. This suggests that the stronger crystallographic textures of the primarily recrystallized grains in the pure iron block the columnar grain growth. This is based on the observance of more twin or low angle boundaries forming between columnar grains and their neighbors. Published data [10] on the correlation between the crystallographic texture and the secondary recrystallization indicated that, when the primary recrystallized texture was strong, certain texture components grew at the expense of others during the course of the secondary recrystallization. A reasonable explanation for this could be that the boundary mobility of grains contributing to this particular component is relatively high. For the directional recrystallization process of pure iron, the specimens underwent primary recrystallization to a fine grain structure when the hot zone was applied to the end of them. For the specimen cold-rolled for 85% of thickness reduction and then directionally annealed at 850 ◦ C, this fine grain structure has recrystallized textures including strong {1 1 0}<1 1 2> and weak {1 1 0}<1 1 0>, {1 1 1}<1 2 1> components, as listed in Table 2. After normal grain growth some grains in the main texture components began to propagate and form a columnar grain along the direction of TG. Grain 3 in Fig. 11 (a) is such an example,
Z.W. Zhang et al. / Materials Science and Engineering A 422 (2006) 241–251
which has the same orientation of {1 1 0}<1 1 2> with the recrystallized texture. With the columnar grains propagating, the tips of the columnar grains grow abnormally at the expense of their small neighboring grains. There is a much higher probability for these columnar grains, which are developed in the main recrystallized textures, to encounter the small grains which have twin or low-angle boundaries, because this kind of columnar grains have the similar orientation to the main recrystallized textures. As a consequence, the columnar grain growth was hindered. 5. Conclusions The effects of the degree of deformation, HZT, specimen drawing velocities and crystallographic textures on the evolution of microstructures have been investigated. The following conclusions were drawn. The columnar grain structure was developed during the secondary recrystallization. Below a critical coarsening temperature, it becomes difficult for the secondary recrystallization to occur and thereby the columnar grain structure cannot form. The degree of deformation is not a key factor in forming columnar grain structure. Directional recrystallization occurs at an optimum specimen drawing velocity at a given HZT and TG. Primary-recrystallized small grains ahead of the hot zone during directional recrystallization have different texture components under different processing conditions. A stronger {1 1 0}<1 1 2> texture was observed in these grains for the specimens cold-rolled for 85% of thickness reduction, while a random texture for those cold-rolled for 70% of thickness reduction. The selective grain growth was by high-angle CSL boundary migration. Twin or low angle boundaries form island grains in the columnar grains due to their low energy and low mobility. The strong crystallographic texture with a {1 1 0}<1 1 2> component in the primarily recrystallized small grains hinders the columnar grain growth due to the higher probability of forming twin or low angle grain boundaries at the tip of the columnar grains. Acknowledgements This work was supported by the Creative Research Foundation for PhD candidates of Jiangsu province and partially by key project of the National Science Foundation (Grant No.:
251
50431030). The authors are grateful to Professor P. Yang and Dr. L. Meng for providing instrument and help with the EBSD work; Professor J. Q. Liu, Professor Y. R. Li and Dr. H. Y. He for some helpful discussions. And the help of Dr. Z. M. Zhang, associate professor J. Kong and Dr. Z. H. Wang and Y Cao with various aspects of the practical work is also gratefully acknowledged. References [1] R.L. Cairns, L.R. Curwick, J.S. Benjamin, Metall. Trans. 6A (1975) 179. [2] R.S. Bellows, E.A. Schwarzkopf, J.K. Tien, Metall. Trans., 19A (1988) 479. [3] B.A. Wilcox, A.H. Clauer, Trans. Metall., Soc AIME 236 (1990) 570. [4] A.W. Godfrey, J.W. Martin, Mater. Sci. Eng. A222 (1997) 91. [5] A. Tetkin, M. Mujahid, J.W. Martin, et al., in: S.D. Antolovich (Ed.), Superalloys, TMS, Warrendale, PA, 1992, p. 457. [6] R.F. Singer, G.H. Gessinger, Metall. Trans. 13A (1982) 1463. [7] J.M. Marsh, J.W. Martin, Mater. Sci. Technol. 7 (1991) 183. [8] A.O. Humphreys, S.W.K. Shaw, J.W. Martin, Mater. Char. 34 (1995) 9. [9] M.S. Greaves, P.S. Bate, W.T. Roberts, et al., Mater. Sci. Technol. 12 (1996) 730. [10] T.S. Chou, H.K.D.H. Bhadeshia, Mater. Sci. Eng. A189 (1994) 229. [11] H.K.D.H. Bhadeshia, Mater. Sci. Eng. A223 (1997) 64. [12] T. Hirano, T. Mawari, Y. Demura, et al., Mater. Sci. Eng. A239-240 (1997) 324. [13] T. Tsujimoto, T. Matsui, T. Suzuki, et al., Intermetallics 9 (2001) 97. [14] I. Baker, J. Li, Acta Mater. 50 (2002) 805. [15] J. Li, S.L. Johns, B.M. Iliescu, Acta Mater. 50 (2002) 4491. [16] S. Ranganathan, Acta Crystallogr. 21 (1966) 197. [17] W.W. Mullins, J. Appl. Phys. 27 (1956) 900. [18] J.E. Taylor, Acta Metal. Mater. 40 (1992) 1475. [19] E.A. Holm, N. Zacharopoulos, D.J. Srolovitz, Acta Mater. 46 (1998) 953. [20] C. Detavernier, S. Rossnagel, C. Noyan, et al., J. Appl. Phys. 94 (2003) 2874. [21] C.V. Thompson, J. Appl. Phys. 58 (1985) 763. [22] J. Li, I. Baker, Mater. Sci. Eng. A392 (2005) 8. [23] R. Carel, C.V. Thompson, H.J. Frost, Acta Mater. 44 (1996) 2479. [24] W.T. Read, W. Shockley, Phys. Rev. 78 (1950) 275. [25] J.E. Burke, D. Turbull, Prog. Metal. Phys. 3 (1952) 220. [26] G. Abbruzzese, K. Lucke, Acta Metal. Mater. 34 (1986) 905. [27] A.D. Rollett, D.J. Srolovitz, M.P. Anderson, Acta Metal Mater. 37 (1989) 1227. [28] F.J. Humphreys, M. Hatherly, Recrystallization and Related Annealing Phenomena, 2nd ed., Elsevier Science Ltd, Oxford, 2004. [29] R. Messina, M. Soucail, T. Baudin, et al., J. Appl. Phys. 84 (1998) 6366. [30] Y. Park, D.-Y. Kim, N.-M. Hwang, J. Appl. Phys. 95 (2004) 5515.
Microstructural evolution of commercial pure iron during directional annealing Z.W. Zhang a,∗ , G.L. Chen a,b , G. Chen a a
Joint Laboratory of Nanostructured Materials and Technology, Nanjing University of Science and Technology, Nanjing 210094, PR China b State key lab for advanced metals and materials, USTB, Beijing 100083, PR China Received 26 January 2006; accepted 3 February 2006
Abstract Microstructural evolution during the directional annealing of the cold-rolled pure iron has been investigated. Columnar grains with the largest aspect ratio of 23.6 were produced in the specimen cold-rolled for 70% of thickness reduction at the hot zone temperature (HZT) of 850 ◦ C and a temperature gradient (TG) of 200 ◦ C/cm with an optimum drawing velocity of 5 m/s. It was found that the columnar grain structure was developed due to abnormal grain growth. There is a critical coarsening temperature, below which the secondary recrystallization does not occur during the annealing; consequently no columnar grain structure can be produced. No obvious evidence shows that the grain size after primary recrystallization has an influence on the microstructural evolution. Due to the strong crystallographic texture of the small grains ahead of the hot zone, the columnar grains have small aspect ratios. Most of the boundaries of the columnar grain are high angle coincidence site lattice (CSL) grain boundaries. The grains with twin or low angle grain boundaries are apt to form island grains in the columnar structure. © 2006 Elsevier B.V. All rights reserved. Keywords: Directional annealing; Columnar-grained structure; Commercial pure iron; Texture
1. Introduction Columnar-grained microstructures, with grain boundaries aligned parallel to the applied tensile stress direction, can enhance creep properties [1–3], improve fatigue resistance [4] and impart crack-stopping behaviors [5]. These structures can often be produced by directional solidification. It may be desirable, however, to use material processed in the solid state. In those cases, directional recrystallization is the possible processing route. Directional recrystallization (zone annealing) is the annealing process whereby the sample is traversed through a furnace (hot zone) containing a steep temperature gradient (with the temperature varying along a direction parallel to the traverse direction) [4,6]. The microstructural evolution during directional recrystallization (DR) depends sensitively on materials and the preparation parameters. In many earlier studies, the DR technique was mainly used on processing superalloys, especially the oxide
∗
Corresponding author. Tel.: +86 025 84315797; fax: +86 025 84315797. E-mail address: [email protected] (Z.W. Zhang).
0921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2006.02.001
dispersion strengthened (ODS) nickel and iron-based superalloys for applications in the aerospace industry [7–13]. For ODS superalloys, the prerequisite for the successful DR is a fine (∼0.2 m) equiaxed matrix grain size [7,8]. In some superalloys, simple static annealing will result in large elongated grains because the initial powder consolidation and rolling or extrusion lead to the alignment of particles which act to pin grain boundaries anisotropically [9–11]. While the principle of such zone annealing has been known for a long time, the process parameters of temperature gradient (TG), hot zone temperature (HZT) and drawing velocity, together with material characteristics, must be closely controlled to produce the desired structure [9]. Humphreys [8] have shown that, in directionally recrystallized MA-6000, drawing velocity can influence the microstructure and the crystallographic texture of MA-6000 extrusions. This can be rationalized using the difference in boundary mobility due to the solute pinning. Recently, Li et al. [15] and Baker and Li [14] researched systematically the effect of processing parameters of the directional annealing on the directional recrysatallization of poly crystalline nickel and cold-rolled copper single crystal. They observed that the processing parameters also influence the number of annealing
242
Z.W. Zhang et al. / Materials Science and Engineering A 422 (2006) 241–251
twins per unit of area and texture evolution, in addition to the microstructure. This work is a part of the study aiming to develop a fundamental understanding of how the initial microstructural parameters and the processing variables determine the microstructural evolution of commercial pure iron during directional annealing. 2. Experimental work The material used in the present work was in the form of bars with an average grain size of 50 m, as determined by the linear intercept method. The microstructure of the as-received materials is shown in Fig. 1. The composition of the commercial pure iron is shown in Table 1. Specimens were cold-rolled to 70% and 85% thickness reduction, with thickness reductions limited to 0.5 mm per pass (<2% of thickness reduction) during the last three passes. The specimens were water-cooled after each pass to prevent heating. The strips after cold rolling for 70% and 85% thickness reduction were 8 and 4.5 mm thick, respectively. And then the strips were cut into bars, which have square crosssections of 8 × 8 mm2 and 4.5 × 4.5 mm2 and are more than 150 mm long. The rolled bars were polished mechanically and then cleaned using acetone. The recrystallization temperatures of the rolled specimens were measured by differential thermal
Fig. 1. Microstructure of as-received commercial pure iron. Table 1 Composition of commercial pure iron in weight percent Element
wt.%
C Si Mn P S Al Ni Cr Cu Ti Fe
0.003 0.025 0.150 0.010 0.010 0.040 / / / / Bal.
Fig. 2. Schematic illustration of the directional annealing furnace. The induction coil is held stationary and the specimen moved from bottom to top or inversely.
analysis (DTA) with a Shimadzu DTA-50 thermal analysis system in nitrogen gas. The specimens were heated from ambient to 1000 ◦ C at a heating rate of 20 ◦ C/min. The measured recrystallization temperatures were 662 and 640 ◦ C for the specimen cold-rolled for 70% and 85% of reduction, respectively. The experimental set-up for the directional annealing is shown in Fig. 2. A 5 mm high single circle induction coil and a Ga-In liquid alloy cooling pool are used. The specimen is held on a transition bar (drawing bar), which is moved through a collet by a stepper motor at speeds ranging from 0.1 m/s to 9000 m/s. The distance between the surface of the Ga–In liquid alloy and the induction coil can be adjusted to change the temperature gradient (TG) ahead of the hot zone. In the present study, the TG was 200 ◦ C/cm and the hot zone temperatures (HZTs) of 700, 750 and 850 ◦ C were used for directional annealing. In several studies of the directional annealing of the nickelbased superalloy INCONEL MA6000 [5], cold-rolled copper single crystals [14] and polycrystalline nickel [15], the specimen was run once for several centimeters at a given temperature but at different moving speeds. In order to distinguish these regions for different speeds, the interval regions of short lengths were run at a very slow rate and a low temperature. As a result, the obtained strip possessed several experimental “runs” with small unrecrystallized intervals. The same technique was employed here. However, in our experimental process, the interval regions were run at a very high rate and a low temperature. With this modification, the same result was obtained; and meanwhile the experimental time was reduced substantially. After directional annealing, the specimens were mechanically polished and etched in 4% nitric acid in ethanol for optical metallography. Grain sizes were measured using the linear intercept method. For specimens with columnar grain structures, grain size measurements were made both parallel and perpendicular to the growth direction. The preferred crystallographic orientations of the as-rolled and the directional-annealed materials were determined by electron backscattered diffraction (EBSD) and analysis of X-ray
Z.W. Zhang et al. / Materials Science and Engineering A 422 (2006) 241–251
reflection pole figures. A Siemens 5000 X-ray diffractometer was employed to measure the textures by recording the 110, 200 and 211 pole figures from longitudinal sections. The Xradiation used was CuK␣ with a wavelength of 0.15406 nm. The EBSD measurements were carried out using a Leo1450 scanning electron microscopy (SEM) equipped with HKL Channel 4.2 software. And TexTools v.3.2, commercial software (ResMat Corp.) was also used for texture analysis. The value of coincidence site lattice (CSL) grain boundary was calculated using Frank-Ranganthan method [16]. 3. Experimental results 3.1. Deformation degree Figs. 3 and 4 show the microstructures of the specimens coldrolled for 70% and 85% of thickness reduction, respectively. The specimens were directionally annealed at 850 ◦ C with various drawing velocities. Obviously, the recrystallized microstructures show different features under different degree of deformation and drawing velocities. At the drawing velocity of 1 m/s, the microstructure consists of large grains, between which there are
243
some small grains (Fig. 3 (a)). Fig. 3(b) shows a columnar grain structure, which was developed at a drawing velocity of 3 m/s. The longitudinal grain size is approximately 6.7 mm and the transverse grain size is 600 m. Increasing the drawing velocity to 5 m/s, the longitudinal and transverse grain sizes are approximately 11.8 mm and 500 m, respectively. The corresponding micrograph is shown in Fig. 3(c). It is worth noting that the boundaries between the large grains grown at the drawing velocities of 3 and 5 m/s are almost parallel to the movement direction of the specimen. As the drawing velocity was increased further to 25 m/s, a small-sized equiaxed grain structure formed with an average grain size of approximately 50 m (Fig. 3(d)). For the degree of deformation of 85% reduction, directional recrystallization occurred at drawing velocity of 3–8 m/s, as shown in Fig. 4. However, the longitudinal grain sizes are smaller and the transverse grain sizes are almost the same comparing with the specimens, which are cold-rolled for reduction of 70%. In addition, we noticed that, for the specimens cold-rolled for 85% reduction, the range of drawing velocity for the occurrence of the directional recrystallization is broader, from 3 to 8 m/s. Fig. 5 shows the profile of the relationship between the grain sizes and the drawing velocities. In order to delineate the grain structure
Fig. 3. Microstructures of commercial pure iron cold-rolled for 70% of thickness reduction and directionally annealed at 850 ◦ C with a temperature gradient of 200 ◦ C/cm and various drawing velocity of (a) 1 m/s, (b) 3 m/s, (c) 5 m/s and (d) 25 m/s. The arrows refer to the direction of columnar grain growth.
244
Z.W. Zhang et al. / Materials Science and Engineering A 422 (2006) 241–251
Fig. 4. Microstructures of commercial pure iron cold-rolled for 85% of thickness reduction and directionally annealed at 850 ◦ C with a temperature gradient of 200 ◦ C/cm and various drawing velocity of (a) 3 m/s, (b) 5 m/s, (c) 8 m/s and (d) 25 m/s. The arrows refer to the direction of columnar grain growth.
more precisely, the grain sizes we used in Fig. 5 were the sizes perpendicular and parallel to the columnar growth direction, separately, instead of the average grain sizes. 3.2. Hot zone temperature (HZT)
Fig. 5. Grain size as a function of drawing velocity for the degree of deformation 70% and 85% reduction at the hot zone temperature of 850 ◦ C and the temperature gradient of 200 ◦ C/cm. When a columnar grain structure is present, the grain size is plotted not as the average grain size but as the grain size perpendicular and parallel to the columnar growth direction.
Fig. 6 shows the grain sizes as a function of drawing velocities for the two HZTs of 700 and 750 ◦ C. Fig. 7 shows the microstructures of the specimens cold-rolled for 85% of thickness reduction. The specimens were directionally annealed at 700 ◦ C with the drawing velocities of 1 m/s and 8 m/s. Smallsized equiaxed grain structures are observed at all the drawing velocities from 1 to 8 m/s. At the drawing velocity of 1 m/s, the average grain size is about 50 m (Fig. 7 (a)). Increasing the drawing velocity, the average grain size decreases slightly. As the drawing velocity was increased further to 8 m/s, a partially recrystallized structure was produced with an average recrystallised grain size of approximately 20 m (Fig. 7(b)). Fig. 8 shows the microstructures of the specimens cold-rolled for reduction of 70%. The specimens directionally annealed at 750 ◦ C with the drawing velocities of 1, 3 and 5 m/s. It can be seen clearly from Fig. 8 (b) that the columnar-like grain struc-
Z.W. Zhang et al. / Materials Science and Engineering A 422 (2006) 241–251
245
ture can be developed only at the drawing velocity of 3 m/s. However, The longitudinal grain size of the elongated grains is much smaller (1.2 mm) than that directionally annealed at 850 ◦ C, although the transverse grain size is almost the same (330 m). At both drawing velocities of 1 and 5 m/s, the grain structures are equiaxed with the average grain size of about 100 m, see Fig. 8 (a) and (c). 3.3. Grain size after primary recrystallization
Fig. 6. Grain size as a function of drawing velocity for the hot zone temperature (HZT) of 700 ◦ C and 750 ◦ C at the temperature gradient of 200 ◦ C/cm. When a columnar grain structure is present, the grain size is plotted not as the average grain size but as the grain size perpendicular and parallel to the columnar growth direction.
To test the effect of grain sizes after primary recrystallization on the columnar grain growth, pure iron specimens cold-rolled for 70% of thickness reduction were isothermally annealed at 650, 700 and 750 ◦ C for 30 min. After isothermally annealed at 650 ◦ C, the specimen had a partially recrystallized structure. The specimens isothermally annealed at 700 ◦ C and 750 ◦ C had completely recrystallized structures with grain size of 50 and 100 m, respectively. These specimens were then directionally annealed at 850 ◦ C and a TG of 200 ◦ C/cm with the drawing velocity of 3 m/s. The resulting microstructures are shown in Fig. 9. The resulting microstructures are very similar to those shown in Fig. 3, which were directionally annealed after cold rolling without isothermal annealing. No obvious evidence indicated that the grain size after primarily recrystallization influences the microstructural evolution. This result supports the suggestion that the columnar grains in the cold-rolled commercial pure iron were produced by the secondary recrystallization. This, perhaps, is not surprising. Under the temperature gradient of 200 ◦ C/cm ahead of the hot zone, at 5 mm ahead of the hot zone the temperature is approximately 700 ◦ C. This is higher than the temperature (662 ◦ C), which allows the commercial pure iron to recrystallize. 3.4. Crystallographic texture
Fig. 7. Microstructures of commercial pure iron cold-rolled for 85% of thickness reduction and directionally annealed at 700 ◦ C and a temperature gradient of 200 ◦ C/cm with drawing velocities of (a) 1 m/s and (b) 8 m/s The direction of annealing was from left to right.
The preferred crystallographic orientations of as cold-rolled texture were determined using data from X-ray reflection pole figures. The orientations of the primary recrystallised small grains ahead of the hot zone and those of the large columnar grains were determined from EBSD attached to a SEM. The textures induced by various processing parameters are summarized in Table 2. Under the cold-rolled condition, the materials had an elongated grain structures. The deformation textures, centered on <1 1 0> fiber with {1 1 1}<1 1 0> and {0 0 1}<1 1 0> components, were developed after reduction of 70% and those, centered on <1 1 0> fiber with {1 1 2}<1 1 0> and {1 1 1}<1 1 0> components, after larger thickness reduction of 85%. When the specimen cold-rolled for 85% of reduction was directionally annealed at 700 ◦ C with the drawing velocity of 1 m/s, recrystallization occurred, resulting in a microstructure with small-sized equiaxed grain, see Fig. 7 (a). Under this condition, this primary recrystallized microstructure had textures, which were different in type from the corresponding cold-rolled textures. There was a tendency for {1 1 2}<1 1 0> component to be depleted. Most primary recrystallized textures were dominated by the near {1 1 1}<1 1 0> component.
246
Z.W. Zhang et al. / Materials Science and Engineering A 422 (2006) 241–251
Fig. 8. Microstructures of commercial pure iron cold-rolled for 70% of thickness reduction and directionally annealed at 750 ◦ C with a temperature gradient of 200 ◦ C/cm and various drawing velocity of (a) 1 m/s, (b) 3 m/s and (c) 5 m/s. The direction of annealing was from left to right.
For the specimens cold-rolled for 85% of reduction and directionally annealed at 850 ◦ C with the drawing velocity of 5 m/s, the recrystallized textures of the small grains ahead of the hot zone are distinctly different from the deformation textures. Most recrystallized textures of the small grains were dominated by the near {1 1 0}<1 1 2> component. Also very weak {1 1 0}<1 1 0>
and {1 1 1}<1 2 1> components are shown in these materials. However, for the specimen cold-rolled for 70% of thickness reduction and directionally annealed with the same processing parameters as those cold-rolled for 85% of thickness reduction, the small grains ahead of the hot zone have a random texture with very weak {1 1 1} <1 2 1> and {1 1 0}<1 1 2> components.
Table 2 Crystallographic textures of commercial pure iron under various processing conditions Degree of deformation (%)
Process variables
Texture
Grain structure
70
As cold-rolled
Elongated grains
85
Rolled and directionally annealed at 850 ◦ C As cold-rolled Rolled and directionally annealed at 700 ◦ C
Stronger components {1 1 1}<1 1 0>, {0 0 1}<1 1 0> Weaker components {1 1 0}<1 1 0>and {1 1 1}<2 1 1> Random texture with very weak {1 1 1} <1 2 1> and {1 1 0}<1 1 2> components Stronger components {1 1 2}<1 1 0> and {1 1 1}<1 1 0>, Weaker components {1 1 0}<1 1 0>, {1 1 1}<2 1 1> <1 1 0>fiber with stronger {1 1 1}<1 1 0> and weaker {1 1 1}<1 2 1>components
Rolled and directionally annealed at 850 ◦ C
Stronger {1 1 0}<1 1 2>, weaker {1 1 0}<1 1 0>and {1 1 1}<1 2 1> components
Columnar grain with high aspect ratio Elongated grains Equiaxed fine grains (Primary recrystallization) Columnar grain with low aspect ratio
For the cold-rolled specimen, texture refers to deformation texture; for those cold-rolled and directionally annealed, textures refer to primarily recrystallized texture of the small grains ahead of the hot zone; Grain structures refer to those after directional grain growth except the as cold-rolled specimens.
Z.W. Zhang et al. / Materials Science and Engineering A 422 (2006) 241–251
247
Fig. 9. Microstructures of commercial pure iron cold-rolled for 70% of reduction and then directionally annealed at 850 ◦ C and a temperature gradient of 200 ◦ C/cm with the drawing velocity of 3 m/s. Before directionally annealed, the cold-rolled specimens were isothermally annealed for 30 min at (a) 650 ◦ C, (b)700 ◦ C and (c) 750 ◦ C. The arrow refers to the direction of columnar grain growth.
We noticed that the specimens cold-rolled for 85% of reduction had the same texture with {1 1 1}<1 1 0> component both when directionally annealed at 850 ◦ C with the drawing velocity of 1000 m/s and directionally annealed at 700 ◦ C with the drawing velocity of 1 m/s, although the former had a smaller grain size of approximately 35 m. Fig. 10 shows the ODF section at Euler angle ϕ2 = 45◦ of the specimens cold-rolled for 85% of reduction and directionally annealed
with the two different processing parameters. This indicates that during the primary recrystallization, the recrystallized textures formed by directional annealing are similar to those by isothermal annealing. Moreover, from comparing the textures of the small equiaxed grains ahead of the hot zone induced by different processing parameters, it is apparent that the recrystallization textures show some dependence on the degree of deformation and heat treatment.
Fig. 10. Section through the ODF at Euler angle 2 = 45◦ of the specimen cold-rolled for 85% of reduction and directionally annealed (a) at 700 ◦ C and a temperature gradient of 200 ◦ C/cm with the drawing velocity of 1 m/s and (b) at 850 ◦ C and the temperature gradient of 200 ◦ C/cm with the drawing velocity of 1000 m/s. The contour levels shown are multiples of random density. The dominant texture components occur close to a line with = 54.7◦ , which corresponds to {1 1 1} lying parallel to the sheet plane. The texture is quite typical of recrystallized ferritic steel plate.
248
Z.W. Zhang et al. / Materials Science and Engineering A 422 (2006) 241–251
Fig. 11 shows the microstructures of the large columnar grains in the specimens cold-rolled for 85% and 70% of thickness reduction. The specimens were directionally annealed at 850 ◦ C with the drawing velocity of 5 m/s. The corresponding orientations of the columnar grains are listed in Table 3. It can be seen from Fig. 11 and Table 3 that the columnar grains after directional annealing have random orientations for both specimens cold-rolled for 85% and 70% of thickness reduction. However, the strong recrystallized texture {1 1 0}<1 1 2> of small grains in the specimen cold-rolled for 85% of thickness reduction produces some columnar grains with the orientation of {1 1 0}<1 1 2> (Grain 3 in Fig. 11 (a)), the recrystallized texture, although of course the statistics of these grains are limited. Under this condition, the columnar grains with smaller aspect ratio were produced. The relationships of the orientation between different columnar grains shown in Fig. 11 are summarized in Table 4. It can be seen that most of the misorientations between the columnar grains are concentrated on high angle coincidence site lattice (CSL) boundaries. Moreover, almost all the island grains (Grain4, 6, 8,12 in Fig. 11 (a) and grain 4, 7 in Fig. 11 (b)) have the twin grain boundaries (60◦ /<1 1 1>)or low angle grain boundaries with respect to the matrix columnar grains. 4. Discussion 4.1. Grain boundary migration velocity
Fig. 11. Microstructures ahead of the hot zone of the specimens cold-rolled for (a) 85% of thickness reduction and (b) 70% of thickness reduction and then directionally annealed at 850 ◦ C with a temperature gradient of 200 ◦ C/cm and the drawing velocity of 5 m/s. The arrows refer to the direction of columnar grain growth.
Microstructural evolution typically involves the movement of grain boundaries within the material. The movement velocity of a grain boundary is given by the well-known rate equation [17–19]. ν = MG where G is the driving force, usually expressed as the weighted mean curvature of the boundary during abnormal grain growth
Table 3 The orientation of the columnar grains in the specimens cold-rolled for thickness reduction of 85% and 70% and directionally annealed at 850 ◦ C and a temperature gradient of 200 ◦ C/cm with the drawing velocity of 5 m/s 85% of thickness reduction Grain no.
1 2 3 4 5 6 7 8 9 10 11 12
70% of thickness reduction
Test Euler angle
(◦ )
(hkl)[uvw]
∅1
∅2
0.4 222.1 325.1 216.9 320.0 217.7 318.8 315.2 58.3 0.4 88.4 0.9
8.3 19.2 45.3 34.2 27.2 33.1 28.5 28.6 37.1 1.2 45.5 27.5
42.6 60.5 91.3 22.9 73.4 21.8 72.4 77.2 28.8 179.2 68.8 1.2
The corresponding microstructures were shown in Fig. 11.
(119)[1-10] (216)[14-1] (101)[1-2-1] (134)[-22-1] (317)[3-2-1] (134)[-22-1] (317)[3-2-1] (317)[3-2-1] (123)[1-53] (001)[-100] (212)[-3-24] (012)[100]
Grain No.
1 2 3 4 5 6 7 8
Test Euler angle (◦ )
(hkl)[uvw]
∅1
∅2
144.8 67.2 261.5 1.2 151.7 201.1 127.9 94.8
30.8 49.2 36.7 46.8 71.3 47.6 34.5 30.5
88.6 35.3 26.1 15.2 64.8 22.8 45.1 74.6
(102)[-231] (233)[0-11] (123)[14-3] (144)[4-10] (943)[-232] (255)[-10 7-3] (112)[-201] (316)[-9-35]
Z.W. Zhang et al. / Materials Science and Engineering A 422 (2006) 241–251
249
Table 4 Misorientation of the columnar grains in the specimens cold-rolled for thickness reduction of 85% and 70% and directionally annealed at 850 ◦ C and a temperature gradient of 200 ◦ C/cm with the drawing velocity of 5 m/s 85% of thickness reduction
70% of thickness reduction
Index1
Index2
Misorientation
GB type
Index1
Index2
Misorientation
GB type
1 2 2 3 4 5 6 7 7 9 10 11
2 3 4 4 5 6 7 8 9 10 11 12
39.1◦ /[0-1-1] 57.7◦ /[1-1-1] 44.8◦ /[-1-4-1] 44.5◦ /[4-4-1] 58.2◦ /[-11-1] 57.0◦ /[-11-1] 58.9◦ /[-11-1] 2.4◦ /[3-3-1] 55.7◦ /[2-1-2] 36.6◦ /[210] 49.8◦ /[5-21] 58.7◦ /[11-1]
9 3 59 229 3 3 3 LA 45 27 87 3
1 2 3 4 5 6 6 7
2 3 5 5 6 7 8 8
45.8◦ /[-121] 44.6◦ /[213] 38.7◦ /[40-1] 58.1◦ /[-1-11] 47.4◦ /[012] 59.2◦ /[-111] 50.6◦ /[-211] 18.1◦ /[30-1]
21 95 161 3 15 3 31 LA
The corresponding microstructures were shown in Fig. 11; Note: Index: the number of the columnar grains marked in the Fig. 11 (a) and (b); LA: low angle boundary.
and M is the grain boundary mobility. Thus, a boundary travels twice as fast when its mobility doubles. Abnormal grain growth is generally assumed to arise from a combination of the inhibition of normal grain growth and the occurrence of a small number of special grains with a low free energy, providing a net driving force for their growth [20]. The resulting texture is claimed to be dependent on the driving force, which selects the “special” grains. Abnormal grain growth implies that one type of grains has exceptional fast-moving boundaries. There are two possible causes: (1) A special driving force G for selected grains (thermodynamics-based selection) and/or (2) a difference in M for different grain boundaries (kinetics-based selection). The typical driving force can be calculated from [21,22] G = Gv + Gb + Gs where Gv is the driving force from energy stored in the deformed matrix in the form of dislocations, Gb is the driving force from decreasing of the grain boundary area and Gs is from the surface energy. As mentioned above, for abnormal grain growth in commercial pure iron during directional annealing, since the columnar grains were produced by secondary recrystallization, the stored energy Gv from cold rolling had been completely consumed by the primary recrystallization. In addition, since the size of the iron bars are 8 mm × 8 mm and 4.5 mm × 4.5 mm in cross-section, the driving force from surface energy can be neglected. Hence, it can be proposed that the only driving force for columnar grain growth is the reduction of grain boundary area. The driving force from reducing of the grain boundary area can be calculated from [23] 1 1 Gb = E − D1 D2 where E is the energy of the grain boundary between the two neighboring grains with radius D1 and D2 , respectively. It is generally accepted that the energy of the grain boundary is a function of misorientation. The famous relationship between them was first derived by Read and Shockley [24] for a simple
tilt boundary based on a dislocation model and is given in the following equation E = E0 θ(A − ln θ) where E is the grain boundary energy, θ is the misorientation angle, E0 and A are constants. The mobility of high angle grain boundaries is temperaturedependent and is often found to obey an Arrhenius equation M = M0 exp(−Q/kT ) where Q is the activation energy for grain boundary migration, k the Boltzmann constant, and M0 is the usual prefactor. As far as the directional annealing is concerned, a mobility gradient implies an effective temperature gradient [25]. In addition, even if under the same high temperature, the differences in mobility may be significant, especially in textured materials [26,27]. The high mobility of the high-angle grain boundaries will cause
Fig. 12. Grain size as a function of drawing velocity for the specimens cold-rolled for 70% of thickness reduction. The specimens were directionally annealed at the hot zone temperature of 750 ◦ C and 850 ◦ C and a temperature gradient of 200 ◦ C/cm with the drawing velocity of 5 m/s. When a columnar grain structure is present, the grain size is plotted not as the average grain size but as the grain size perpendicular and parallel to the columnar growth direction.
250
Z.W. Zhang et al. / Materials Science and Engineering A 422 (2006) 241–251
some “special” grains grow rapidly by a kinetics-based selection mechanism during directional annealing. 4.2. Effect of processing variables From the studies mentioned above, it is clearly that the microstructure evolution strongly depends on the processing parameters. Fig. 12 shows the effect of the HZT on the grain size and shape. At the same TG and the drawing velocities, the higher HZT (850 ◦ C) benefits the forming of columnar grain structures. This suggests that, for the secondary recrystallization, there is a critical temperature, which is also called temperature of grain coarsening. When the annealing temperature is below this temperature the abnormal grain growth cannot occur, so the large columnar grain structure cannot be developed, see Figs. 3 and 8. The aspect ratios of the columnar grains under various processing parameters were shown in Fig. 13. The degree of deformation is not the key factor in determining the columnar grain growth. The specimens cold-rolled for 85% of reduction have more stored energy than those cold-rolled for 70% of reduction, however, the aspect ratio of columnar grains in these materials is smaller. This, perhaps, is not surprising, because the columnar grains were produced not by primary recrystallization but by secondary recrystallization. During primary recrystallization the stored energy drives the nucleation and normal grain growth. After the primary recrystallization, the store energy is depleted. At least, the residual strain energy is not the key driving force for the directional grain growth. Drawing velocity has an important influence on the aspect ratio of columnar grains. There is an optimum range of drawing velocities for columnar grain growth. The interpretation of the influence of drawing velocities on the microstructural evolution during the directional annealing requires consideration of the two distinct velocities involved in columnar grain growth [19]: the drawing velocity and the grain boundary movement. Holm et al. [19] simulated the microstructural evolution during zone annealing with different hot zone moving velocity (expressed as
Fig. 13. Aspect ratio of grains as a function of drawing velocity under various processing parameters. When an equiaxed grain structure is present, the aspect ratio equals to 1.
drawing velocity in the present paper). At small drawing velocities, the grain boundaries in the region ahead of the hot zone move faster than the drawing velocity. Thus, the fastest evolution is due to normal grain growth in this region, with a little extra growth while the hot zone is present. This leads to a largesized, equiaxed microstructure throughout the sample, as shown in Fig. 3 (a). When the drawing velocity is equal to the velocity of the boundaries in the hot zone, columnar grain growth occurs. If the drawing velocity is faster than that of the boundaries, both the length and width of grains are approximately equal to their normal grain size because the hot zone sweeps through the specimen too quickly to influence the boundaries. 4.3. Effect of crystallographic texture The columnar grains, which contact their neighbors by highangle CSL boundaries, propagate continuously and preferentially along the direction of TG, as shown in Fig. 11 and Table 4. The selective high-angle CSL boundary migration can be explained by the higher impurity concentration in the general high-angle grain boundaries than in high-angle CSL boundaries, which prevented the movement of general high-angle grain boundaries. It has been known [28] that the effects of impurities on boundary mobility are dependent on the crystallography of the boundary. The special boundaries, i.e. those which are close to a coincidence relationship, are less susceptible to effects of impurities than random boundaries. According to the findings reported by Messina et al. [29] and Park et al. [30], 3 and5 boundaries are expected to have low energy and low mobility. It is easy for these kinds of boundaries to form island grains, growing out into neighboring columnar grains. The similar result was found in this study. As noted above, directional secondary recrystallization of heavily cold-rolled commercial pure iron produced columnar grains with smaller aspect ratio than those in specimens coldrolled for 70% of thickness reduction. This suggests that the stronger crystallographic textures of the primarily recrystallized grains in the pure iron block the columnar grain growth. This is based on the observance of more twin or low angle boundaries forming between columnar grains and their neighbors. Published data [10] on the correlation between the crystallographic texture and the secondary recrystallization indicated that, when the primary recrystallized texture was strong, certain texture components grew at the expense of others during the course of the secondary recrystallization. A reasonable explanation for this could be that the boundary mobility of grains contributing to this particular component is relatively high. For the directional recrystallization process of pure iron, the specimens underwent primary recrystallization to a fine grain structure when the hot zone was applied to the end of them. For the specimen cold-rolled for 85% of thickness reduction and then directionally annealed at 850 ◦ C, this fine grain structure has recrystallized textures including strong {1 1 0}<1 1 2> and weak {1 1 0}<1 1 0>, {1 1 1}<1 2 1> components, as listed in Table 2. After normal grain growth some grains in the main texture components began to propagate and form a columnar grain along the direction of TG. Grain 3 in Fig. 11 (a) is such an example,
Z.W. Zhang et al. / Materials Science and Engineering A 422 (2006) 241–251
which has the same orientation of {1 1 0}<1 1 2> with the recrystallized texture. With the columnar grains propagating, the tips of the columnar grains grow abnormally at the expense of their small neighboring grains. There is a much higher probability for these columnar grains, which are developed in the main recrystallized textures, to encounter the small grains which have twin or low-angle boundaries, because this kind of columnar grains have the similar orientation to the main recrystallized textures. As a consequence, the columnar grain growth was hindered. 5. Conclusions The effects of the degree of deformation, HZT, specimen drawing velocities and crystallographic textures on the evolution of microstructures have been investigated. The following conclusions were drawn. The columnar grain structure was developed during the secondary recrystallization. Below a critical coarsening temperature, it becomes difficult for the secondary recrystallization to occur and thereby the columnar grain structure cannot form. The degree of deformation is not a key factor in forming columnar grain structure. Directional recrystallization occurs at an optimum specimen drawing velocity at a given HZT and TG. Primary-recrystallized small grains ahead of the hot zone during directional recrystallization have different texture components under different processing conditions. A stronger {1 1 0}<1 1 2> texture was observed in these grains for the specimens cold-rolled for 85% of thickness reduction, while a random texture for those cold-rolled for 70% of thickness reduction. The selective grain growth was by high-angle CSL boundary migration. Twin or low angle boundaries form island grains in the columnar grains due to their low energy and low mobility. The strong crystallographic texture with a {1 1 0}<1 1 2> component in the primarily recrystallized small grains hinders the columnar grain growth due to the higher probability of forming twin or low angle grain boundaries at the tip of the columnar grains. Acknowledgements This work was supported by the Creative Research Foundation for PhD candidates of Jiangsu province and partially by key project of the National Science Foundation (Grant No.:
251
50431030). The authors are grateful to Professor P. Yang and Dr. L. Meng for providing instrument and help with the EBSD work; Professor J. Q. Liu, Professor Y. R. Li and Dr. H. Y. He for some helpful discussions. And the help of Dr. Z. M. Zhang, associate professor J. Kong and Dr. Z. H. Wang and Y Cao with various aspects of the practical work is also gratefully acknowledged. References [1] R.L. Cairns, L.R. Curwick, J.S. Benjamin, Metall. Trans. 6A (1975) 179. [2] R.S. Bellows, E.A. Schwarzkopf, J.K. Tien, Metall. Trans., 19A (1988) 479. [3] B.A. Wilcox, A.H. Clauer, Trans. Metall., Soc AIME 236 (1990) 570. [4] A.W. Godfrey, J.W. Martin, Mater. Sci. Eng. A222 (1997) 91. [5] A. Tetkin, M. Mujahid, J.W. Martin, et al., in: S.D. Antolovich (Ed.), Superalloys, TMS, Warrendale, PA, 1992, p. 457. [6] R.F. Singer, G.H. Gessinger, Metall. Trans. 13A (1982) 1463. [7] J.M. Marsh, J.W. Martin, Mater. Sci. Technol. 7 (1991) 183. [8] A.O. Humphreys, S.W.K. Shaw, J.W. Martin, Mater. Char. 34 (1995) 9. [9] M.S. Greaves, P.S. Bate, W.T. Roberts, et al., Mater. Sci. Technol. 12 (1996) 730. [10] T.S. Chou, H.K.D.H. Bhadeshia, Mater. Sci. Eng. A189 (1994) 229. [11] H.K.D.H. Bhadeshia, Mater. Sci. Eng. A223 (1997) 64. [12] T. Hirano, T. Mawari, Y. Demura, et al., Mater. Sci. Eng. A239-240 (1997) 324. [13] T. Tsujimoto, T. Matsui, T. Suzuki, et al., Intermetallics 9 (2001) 97. [14] I. Baker, J. Li, Acta Mater. 50 (2002) 805. [15] J. Li, S.L. Johns, B.M. Iliescu, Acta Mater. 50 (2002) 4491. [16] S. Ranganathan, Acta Crystallogr. 21 (1966) 197. [17] W.W. Mullins, J. Appl. Phys. 27 (1956) 900. [18] J.E. Taylor, Acta Metal. Mater. 40 (1992) 1475. [19] E.A. Holm, N. Zacharopoulos, D.J. Srolovitz, Acta Mater. 46 (1998) 953. [20] C. Detavernier, S. Rossnagel, C. Noyan, et al., J. Appl. Phys. 94 (2003) 2874. [21] C.V. Thompson, J. Appl. Phys. 58 (1985) 763. [22] J. Li, I. Baker, Mater. Sci. Eng. A392 (2005) 8. [23] R. Carel, C.V. Thompson, H.J. Frost, Acta Mater. 44 (1996) 2479. [24] W.T. Read, W. Shockley, Phys. Rev. 78 (1950) 275. [25] J.E. Burke, D. Turbull, Prog. Metal. Phys. 3 (1952) 220. [26] G. Abbruzzese, K. Lucke, Acta Metal. Mater. 34 (1986) 905. [27] A.D. Rollett, D.J. Srolovitz, M.P. Anderson, Acta Metal Mater. 37 (1989) 1227. [28] F.J. Humphreys, M. Hatherly, Recrystallization and Related Annealing Phenomena, 2nd ed., Elsevier Science Ltd, Oxford, 2004. [29] R. Messina, M. Soucail, T. Baudin, et al., J. Appl. Phys. 84 (1998) 6366. [30] Y. Park, D.-Y. Kim, N.-M. Hwang, J. Appl. Phys. 95 (2004) 5515.