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Effect of crystallographic texture on the isothermal beta grain-growth kinetics of Ti–6Al–4V
Materials Science and Engineering A332 (2002) 343– 350 www.elsevier.com/locate/msea
Effect of crystallographic texture on the isothermal beta grain-growth kinetics of Ti–6Al–4V O.M. Ivasishin a,*, S.V. Shevchenko a, S.L. Semiatin b b
a Institute for Metal Physics, 36 Vernadsky Street, 03142 Kie6, Ukraine Air Force Research Laboratory, Materials and Mfg Directorate, AFRL/MLLM, Wright-Patterson Air Force Base, OH 45433 -7817, USA
Received 22 June 2001; received in revised form 23 July 2001
Abstract The effect of crystallographic texture on the kinetics of grain growth during isothermal beta annealing of Ti– 6Al–4V was established. For this purpose, samples were subjected to a thermomechanical process to produce material with a uniform, fine equiaxed-alpha microstructure with either a strong or weak (110) beta-phase texture. Grain growth measurements revealed that the classical isothermal grain-growth relation was incapable of describing the observations. Specifically, the grain-growth exponent n was found to have a strong dependence on temperature. In addition, the phenomenon of discontinuous grain growth, or a stagnation in the growth behavior at short times, was documented and ascribed to changes in texture during the grain-growth process. Published by Elsevier Science B.V. Keywords: Annealing; Titanium; Grain growth; Texture
1. Introduction The beta-grain size is an important factor in determining the mechanical properties of beta heat-treated titanium alloys [1]. The average grain size that characterizes the microstructure of these alloys after heat treatment depends strongly on grain-growth kinetics which are a complex function of a number of factors. For the simplest case of normal grain growth, isothermal grain-growth kinetics can be expressed by the following equation: D n −D n0 =Ktexp
−Q , RT
(1)
in which Q is the activation energy, D and D0 denote the initial and final grain sizes, t is the annealing time, R is the gas constant, T is the absolute temperature, n is the grain growth exponent, and K represents the rate constant. The exponential part of the Eq. (1) comes from the temperature dependence of the boundary mo* Corresponding author. Tel.: + 380-44-444-2210; fax: + 380-44444-0120. E-mail address: [email protected] (O.M. Ivasishin). 0921-5093/02/$ - see front matter. Published by Elsevier Science B.V. PII: S 0 9 2 1 - 5 0 9 3 ( 0 1 ) 0 1 7 5 5 - 5
bility [2] which is closely related to local diffusion processes near or in the moving interface and is usually given by an Arrhenius expression, viz.: M(T)=M0exp
−Q , RT
(2)
where M0 is the pre-exponential factor. Eq. (1) is often utilized to determine n and Q from experimental data on isothermal grain growth. Assuming that the initial grain size D0 is small, the graingrowth exponent n is equal to the inverse of the slope of a ln D versus ln t plot. Similarly, Q can be established from the slopes of plots of ln D versus ln (1/T) at fixed times, i.e.: Q d(ln D) =− . nR d(ln(1/T)
(3)
After determining n and Q, the value of K can be determined using a fitting procedure with the experimental data. The beta grain-growth kinetics for titanium alloys heat treated using both conventional, long-time (furnace) and short-time (salt pot, induction) methods have been the subject of many studies [3–8]. A careful
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Fig. 1. Grain growth data for IMI 685 illustrating (a) normal growth (n = 3.03); and (b) slow growth (n= 11.0) [3]. Fig. 4. Optical microstructure along three orthogonal planes of Ti– 6Al– 4V plate processed as in Fig. 2. The reduction was 87% in four passes.
Fig. 2. Schematic illustration of the thermomechanical processing employed to obtain uniform fine microstructure.
Fig. 3. Optical microstructure of Ti – 6Al–4V program alloy after solution treatment at 1100 °C for 0.5 h followed by forced-convection cooling.
analysis of the results of these investigations reveals deviations from classical (normal) grain-growth behavior whose source has yet to be definitively identified. For instance, Fig. 1 shows an example of noticeably different behaviors for two lots of IMI 685 of the same nominal chemical composition [3]. The phenomenon of ‘discontinuous’ beta grain growth was also reported for Ti –6Al –4V in [3]. Possible explanations for discontinu-
ous grain growth, including solute drag, particle pinning, and crystallographic texture, were proposed; texture was hypothesized to be the most likely reason. By contrast, the opposite conclusion, i.e. that texture does not play an important role in grain growth during beta annealing of Ti–6Al –4V, was reached in [6,7]. However, the textures of the program materials were not well controlled in these latter efforts. In a subsequent investigation [9], the influence of crystallographic texture on the beta grain-growth kinetics of Ti –6Al –4V with different textures during continuous heat treatment was studied. It was assumed that the material coefficients n, Q, and K in the nonisothermal version of Eq. (1) were independent of temperature. Observations of varying grain-growth kinetics were explained in terms of the effect of texture on grain boundary surface energy/mobility. The purpose of the current investigation was to clarify the effect of texture on beta grain growth under isothermal heat treatment conditions. For this purpose, detailed, carefully-controlled experiments were conducted on two lots of specially processed Ti–6Al –4V that differed only in their crystallographic texture.
2. Materials The program material was Ti–6Al – 4V received as hot-rolled plate whose thickness was 16 mm. Its measured composition (in wt.%) was 6.05 Al, 4.40 V, 0.15 Fe, 0.12 Si, 0.02 Mo, 0.16 O, 0.03 N, 0.007 H, balance titanium. The Ti –6Al –4V plate was further processed to obtain two lots of material which met two criteria—(1) both lots should be chemically and microstructurally equivalent, but have noticeably different crystallo-
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graphic textures; and (2) the microstructures of both lots should be equiaxed and as fine and uniform as possible. These requirements can be met via a thermomechanical processing method [10] consisting of the rolling to relatively high reductions of beta-solutionedand-forced-convection-cooled material followed by recrystallization annealing in the a +b phase field (Fig. 2). By this means, one can obtain various textures by changing the temperature of deformation, while the annealing temperature controls the volume fraction and size of the equiaxed, primary-alpha phase [10]. In this approach, prior beta solutioning is an important step. Its objective is to eliminate or minimize the influence of the as-received microstructure and texture on the resultant features of the subsequently-rolled product inasmuch as the primary-alpha phase is fully dissolved, and a fine lamellar-type microstructure is formed (Fig. 3). Preliminary experiments verified that rolling temperature and reduction are the two major parameters controlling the dispersion and uniformity of the final
345
microstructure. The most uniform, equiaxed microstructure with an average alpha grain size of 2–3 mm was achieved by employing rolling reductions of not less than 80% at 850 °C (Fig. 4). With regard to texture, pole-figure analysis indicated that the as-received material was only weakly textured. Furthermore, beta annealing prior to rolling eliminated the initial texture and resulted in the formation of a near-random texture through the b a phase transformation. Rolling at 850 °C to a total reduction in thickness of 80% and a low reduction per pass (10%) led to the formation of a basal type alpha-phase texture. On the other hand, rolling at the same temperature to a total reduction of 87% using a high reduction per pass led to the formation of a basal-transverse texture of the alpha phase (Fig. 5). Based on the above observations, two lots of material having equivalent chemistry and microstructure, but different crystallographic textures of the alpha phase (specifically, basal and basal-transverse) were fabri-
Fig. 5. Alpha-phase (0002) pole figures for (a) lot PM; and (b) lot SM.
Fig. 6. Beta-phase (110) pole figures for (a) lot PM; and (b) lot SM.
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Fig. 7. Optical microstructures of Ti – 6Al– 4V samples after isothermal annealing: (a) PM at 1030 °C for 30 min; (b) SM at 1030 °C for 30 min; (c) PM at 1100 °C for 15 min; and (d) SM at 1100 °C for 15 min.
cated. The first was processed by rolling at 850 °C from 16 to 5 mm in eight rolling passes; the second was rolled at the same temperature from 16 to 2 mm in four rolling passes. Due to the difference in the final thickness, the materials are hereafter referred to as PM (plate material) and SM (sheet material), respectively. The textures of the beta phase of the PM and SM materials, which would impact the subsequent beta grain growth due to the epitaxial growth of beta phase during the dissolution of alpha early in the annealing process, are shown in Fig. 6. The (110) pole figures for PM and SM were generally of the same type, but were much sharper in SM due to differences in the degree of continuous dynamic recrystallization (during rolling) and static discontinuous recrystallization (during interpass reheating) imparted by the thermomechanical processing of the two lots of material. Such textures, however, are generally typical of rolled bcc metals.
3. Experimental procedures Isothermal beta annealing of the PM and SM lots of Ti – 6Al –4V was conducted to determine the effect of texture on beta grain growth. To this end, samples of each material were heat treated in a vacuum furnace at temperatures of 1000, 1030, 1060, 1100, 1130, 1160, 1200, 1230, and 1260 °C for times ranging from 5 to 90 min. The relatively small increment in temperature (i.e. 30 °C) used in these experiments was chosen to provide sufficient data to ensure a very accurate determination of the grain-growth activation energy Q and graingrowth exponent n in Eq. (1) within the given temperature range. After isothermal exposure, each specimen was cooled rapidly by turning off the power to the water-cooled furnace. This yielded an initial sample cooling rate of not less than 50 K s − 1, which was sufficient to stop beta grain growth. At temperatures
O.M. I6asishin et al. / Materials Science and Engineering A332 (2002) 343–350
below the beta transus, the sample cooling rate was somewhat lower, thereby promoting diffusional decomposition of the beta phase. By this means, the beta grain boundaries were decorated with a thin layer of alpha that was beneficial in determining the beta grain size during subsequent optical metallography. Average grain sizes were determined by using the linear intercept method; grains close to the surface were not taken into account to avoid the pinning effect of the free surface. For both materials, the standard deviation was approximately 17–20% of the average grain size with a confidence interval (CI) of 95%.
4. Results and discussion Typical microstructures of the SM and PM materials after identical isothermal annealing treatments are compared in Fig. 7. In all cases, a uniform equiaxed beta grain structure was developed. Some quantitative grain-size measurements are summarized in Fig. 8; the average beta
Fig. 8. Average beta grain size as a function of annealing time at (a) 1000 °C; or (b) 1200 °C.
347
grain size as a function of annealing time at 1000 or 1200 °C is shown for the PM and SM materials. By and large, the maximum grain sizes were approximately 1.0–1.2 mm for PM and 0.3– 0.4 mm for SM, or values substantially smaller than the smallest cross section dimension (= the sheet per plate thickness) in each case. Grain size measurements for all of the annealing temperatures are shown in Fig. 9a (PM) and Fig. 9b (SM). An inspection of the data in Figs. 8 and 9 reveals that the PM material exhibited much faster beta grain growth than the SM material. Equally important, the experimental data for SM had marked indications of discontinuous beta grain growth characterized by a significant period of stagnation when the grain size was about 250– 300 mm. On the other hand, the PM material exhibited rather typical isothermal grain-growth behavior. Plots of ln D versus ln t (Figs. 10 and 11) were made for both materials to determine the value of the graingrowth exponent n and thus the basis for applying Eq. (1) to the experimental data. However, these graphs showed that n did not remain constant over the range of temperatures investigated in the present work. Moreover, the ln D versus ln t graphs were quite different for PM and SM. For the PM material (Fig. 10), grain growth was characterized by a constant value of n for each annealing temperature; the specific value, however, was temperature dependent (Fig. 12). A sigmoidal dependence n(T) was found, thus indicating slower grain growth at higher temperatures when the grain size was large. Despite the fact that the ln D versus ln(1/T) graph is linear for the PM material (Fig. 13), the evaluation of the activation energy Q could not be done because of the sigmoidal n(T) dependence. For the SM material, the n(T) dependence was even more complicated (Fig. 11). At low annealing temperatures, the slope of the ln D versus ln t plots (Fig. 11a) yielded an average n of approximately 8.9, or a value indicative of very slow grain growth. At temperatures above 1100 °C (Fig. 11b), the grain-growth rate increased with both temperature and time. The stage of grain-growth stagnation observed at these temperatures when the average beta grain size was about 250– 300 mm changed to relatively fast growth at longer times. The higher the temperature, the shorter was the exposure necessary for fast growth to start. Using data for 15–60 min heat treatments, an n value of 3.9 was derived for the higher-temperature behavior. With such an n value, the activation energy Q for SM was estimated as 227 kJ mol − 1 from the corresponding ln D versus ln(1/T) graph (Fig. 14). The observations from the present investigation show that materials with different initial textures exhibit noticeably different grain-growth kinetics during isothermal heat treatments. The data also revealed indications of discontinuous growth characterized by alternating stages of relatively fast (normal) grain growth and relatively slow growth (Fig. 9). This phenomenon was
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Fig. 9. Average beta grain size as a function of annealing time for (a) PM; or (b) SM.
Fig. 10. ln D versus ln t dependence for PM material at various temperatures.
apparent in SM at all temperatures. An initial period of fast growth during the first 5 min to a grain size of approximately 250 mm was followed by very slow growth and then a second stage of fast growth. The time interval of slow growth decreased as the annealing temperature
was raised. The annealing temperature also controlled the final grain size that was achieved by the end of a 1-h exposure; the final grain sizes were 400, 600, 750, and 900 mm at 1130, 1160, 1200, and 1230 °C, respectively. In the PM material, discontinuous grain growth was not as
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evident for the exposure times employed in the present work. An increase in grain growth rate at longer times was observed only at high temperatures; at these temperatures, the initial value of n was approximately 5.6, which itself corresponds to relatively slow growth. The present observations agree well with previous, less detailed measurements of discontinuous grain growth for
Fig. 13. ln D versus ln(1/T) for PM material at t= 15 min.
Fig. 14. ln D versus ln(1/T) for SM material at t = 15 min.
Fig. 11. ln D versus ln t dependence for SM material at various temperatures.
IMI685 and Ti–6Al –4V [3]. The experimental data can also be rationalized on the basis of previous analytical modeling [11] that suggests the possibility of discontinuous grain growth in initially strongly textured materials. This modeling work revealed that a relatively strong initial texture could be the most likely reason for very slow initial grain growth and much quicker growth after a reduction in the sharpness of the texture at long annealing times.
5. Conclusions
Fig. 12. Temperature dependence of the grain growth exponent n for PM material.
A series of carefully controlled, isothermal beta annealing trials were conducted on samples of Ti– 6Al – 4V processed to yield identical, fine-grained starting microstructures but different crystallographic textures. Measurements of beta grain growth kinetics led to the following conclusions: (1) Beta-grain growth under isothermal annealing conditions is strongly affected by texture in materials which are otherwise identical in
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terms of chemical composition and microstructure. (2) The kinetics of beta grain growth in titanium alloys can not be generally described in terms of a constant (temperature-independent) grain-growth exponent and activation energy. (3) Strict control of initial texture is required in thermomechanical processing of titanium alloys to ensure the desired level of those mechanical properties which are affected by the beta grain size.
Acknowledgements The present work was supported by the Air Force Office of Scientific Research (AFOSR) and the AFOSR European Office of Aerospace Research and Development (AFOSR/EOARD) within the framework of STCU Partner Project P-041. The encouragement of the AFOSR program managers (Dr R.S. Fredell and Dr C.S. Hartley) is greatly appreciated.
References [1] V.N. Gridnev, O.M. Ivasishin, S.P. Oshkadjorov, Physical Principles of Rapid Heat Treatment of Titanium Alloys (in Russian), Naukova Dumka, Kyiv, 1986. [2] P. Gordon, Energetics in Metallurgical Phenomena, vol. 1, Gordon and Breach, New York, 1962, p. 207. [3] S.P. Fox, in: F.H. Froes, I.L. Caplan (Eds.), Titanium’92, Science and Technology, TMS, Warrendale, PA, 1992, pp. 769 –776. [4] S.L. Semiatin, J. Soper, I.M. Sukonnik, Acta Mater. 44 (1996) 1979. [5] S.L. Semiatin, J. Soper, I.M. Sukonnik, Scr. Metall. et Mater. 30 (1994) 951. [6] F.G. Gil, P. Tarin, J.A. Planell, in: F.H. Froes, I.L. Caplan (Eds.), Titanium’92, Science and Technology, TMS, Warrendale, PA, 1992, pp. 777 – 784. [7] F.G. Gil, J.A. Planell, Mater. Sci. Eng., A 283 (2000) 17. [8] O.M. Ivasishin, P.E. Markovsky, L. Wagner, G. Luetjering, Titanium 1990: Products and Applications, Titanium Development Association, Dayton, OH, 1990, pp. 99 – 110. [9] S.L. Semiatin, P.N. Fagin, M.G. Glavicic, I.M. Sukonnik, O.M. Ivasishin, Mater. Sci. Eng., A 299 (2001) 225. [10] G. Luetjering, M. Peters, Mechanical Properties of a Titanium Blading Alloy, Report CS-2933, Electric Power Research Institute, Palo Alto, CA, 1983. [11] H. Eichelkraut, G. Abbruzzese, K. Lucke, Acta Metall. 36 (1988) 55.
Effect of crystallographic texture on the isothermal beta grain-growth kinetics of Ti–6Al–4V O.M. Ivasishin a,*, S.V. Shevchenko a, S.L. Semiatin b b
a Institute for Metal Physics, 36 Vernadsky Street, 03142 Kie6, Ukraine Air Force Research Laboratory, Materials and Mfg Directorate, AFRL/MLLM, Wright-Patterson Air Force Base, OH 45433 -7817, USA
Received 22 June 2001; received in revised form 23 July 2001
Abstract The effect of crystallographic texture on the kinetics of grain growth during isothermal beta annealing of Ti– 6Al–4V was established. For this purpose, samples were subjected to a thermomechanical process to produce material with a uniform, fine equiaxed-alpha microstructure with either a strong or weak (110) beta-phase texture. Grain growth measurements revealed that the classical isothermal grain-growth relation was incapable of describing the observations. Specifically, the grain-growth exponent n was found to have a strong dependence on temperature. In addition, the phenomenon of discontinuous grain growth, or a stagnation in the growth behavior at short times, was documented and ascribed to changes in texture during the grain-growth process. Published by Elsevier Science B.V. Keywords: Annealing; Titanium; Grain growth; Texture
1. Introduction The beta-grain size is an important factor in determining the mechanical properties of beta heat-treated titanium alloys [1]. The average grain size that characterizes the microstructure of these alloys after heat treatment depends strongly on grain-growth kinetics which are a complex function of a number of factors. For the simplest case of normal grain growth, isothermal grain-growth kinetics can be expressed by the following equation: D n −D n0 =Ktexp
−Q , RT
(1)
in which Q is the activation energy, D and D0 denote the initial and final grain sizes, t is the annealing time, R is the gas constant, T is the absolute temperature, n is the grain growth exponent, and K represents the rate constant. The exponential part of the Eq. (1) comes from the temperature dependence of the boundary mo* Corresponding author. Tel.: + 380-44-444-2210; fax: + 380-44444-0120. E-mail address: [email protected] (O.M. Ivasishin). 0921-5093/02/$ - see front matter. Published by Elsevier Science B.V. PII: S 0 9 2 1 - 5 0 9 3 ( 0 1 ) 0 1 7 5 5 - 5
bility [2] which is closely related to local diffusion processes near or in the moving interface and is usually given by an Arrhenius expression, viz.: M(T)=M0exp
−Q , RT
(2)
where M0 is the pre-exponential factor. Eq. (1) is often utilized to determine n and Q from experimental data on isothermal grain growth. Assuming that the initial grain size D0 is small, the graingrowth exponent n is equal to the inverse of the slope of a ln D versus ln t plot. Similarly, Q can be established from the slopes of plots of ln D versus ln (1/T) at fixed times, i.e.: Q d(ln D) =− . nR d(ln(1/T)
(3)
After determining n and Q, the value of K can be determined using a fitting procedure with the experimental data. The beta grain-growth kinetics for titanium alloys heat treated using both conventional, long-time (furnace) and short-time (salt pot, induction) methods have been the subject of many studies [3–8]. A careful
344
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Fig. 1. Grain growth data for IMI 685 illustrating (a) normal growth (n = 3.03); and (b) slow growth (n= 11.0) [3]. Fig. 4. Optical microstructure along three orthogonal planes of Ti– 6Al– 4V plate processed as in Fig. 2. The reduction was 87% in four passes.
Fig. 2. Schematic illustration of the thermomechanical processing employed to obtain uniform fine microstructure.
Fig. 3. Optical microstructure of Ti – 6Al–4V program alloy after solution treatment at 1100 °C for 0.5 h followed by forced-convection cooling.
analysis of the results of these investigations reveals deviations from classical (normal) grain-growth behavior whose source has yet to be definitively identified. For instance, Fig. 1 shows an example of noticeably different behaviors for two lots of IMI 685 of the same nominal chemical composition [3]. The phenomenon of ‘discontinuous’ beta grain growth was also reported for Ti –6Al –4V in [3]. Possible explanations for discontinu-
ous grain growth, including solute drag, particle pinning, and crystallographic texture, were proposed; texture was hypothesized to be the most likely reason. By contrast, the opposite conclusion, i.e. that texture does not play an important role in grain growth during beta annealing of Ti–6Al –4V, was reached in [6,7]. However, the textures of the program materials were not well controlled in these latter efforts. In a subsequent investigation [9], the influence of crystallographic texture on the beta grain-growth kinetics of Ti –6Al –4V with different textures during continuous heat treatment was studied. It was assumed that the material coefficients n, Q, and K in the nonisothermal version of Eq. (1) were independent of temperature. Observations of varying grain-growth kinetics were explained in terms of the effect of texture on grain boundary surface energy/mobility. The purpose of the current investigation was to clarify the effect of texture on beta grain growth under isothermal heat treatment conditions. For this purpose, detailed, carefully-controlled experiments were conducted on two lots of specially processed Ti–6Al –4V that differed only in their crystallographic texture.
2. Materials The program material was Ti–6Al – 4V received as hot-rolled plate whose thickness was 16 mm. Its measured composition (in wt.%) was 6.05 Al, 4.40 V, 0.15 Fe, 0.12 Si, 0.02 Mo, 0.16 O, 0.03 N, 0.007 H, balance titanium. The Ti –6Al –4V plate was further processed to obtain two lots of material which met two criteria—(1) both lots should be chemically and microstructurally equivalent, but have noticeably different crystallo-
O.M. I6asishin et al. / Materials Science and Engineering A332 (2002) 343–350
graphic textures; and (2) the microstructures of both lots should be equiaxed and as fine and uniform as possible. These requirements can be met via a thermomechanical processing method [10] consisting of the rolling to relatively high reductions of beta-solutionedand-forced-convection-cooled material followed by recrystallization annealing in the a +b phase field (Fig. 2). By this means, one can obtain various textures by changing the temperature of deformation, while the annealing temperature controls the volume fraction and size of the equiaxed, primary-alpha phase [10]. In this approach, prior beta solutioning is an important step. Its objective is to eliminate or minimize the influence of the as-received microstructure and texture on the resultant features of the subsequently-rolled product inasmuch as the primary-alpha phase is fully dissolved, and a fine lamellar-type microstructure is formed (Fig. 3). Preliminary experiments verified that rolling temperature and reduction are the two major parameters controlling the dispersion and uniformity of the final
345
microstructure. The most uniform, equiaxed microstructure with an average alpha grain size of 2–3 mm was achieved by employing rolling reductions of not less than 80% at 850 °C (Fig. 4). With regard to texture, pole-figure analysis indicated that the as-received material was only weakly textured. Furthermore, beta annealing prior to rolling eliminated the initial texture and resulted in the formation of a near-random texture through the b a phase transformation. Rolling at 850 °C to a total reduction in thickness of 80% and a low reduction per pass (10%) led to the formation of a basal type alpha-phase texture. On the other hand, rolling at the same temperature to a total reduction of 87% using a high reduction per pass led to the formation of a basal-transverse texture of the alpha phase (Fig. 5). Based on the above observations, two lots of material having equivalent chemistry and microstructure, but different crystallographic textures of the alpha phase (specifically, basal and basal-transverse) were fabri-
Fig. 5. Alpha-phase (0002) pole figures for (a) lot PM; and (b) lot SM.
Fig. 6. Beta-phase (110) pole figures for (a) lot PM; and (b) lot SM.
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Fig. 7. Optical microstructures of Ti – 6Al– 4V samples after isothermal annealing: (a) PM at 1030 °C for 30 min; (b) SM at 1030 °C for 30 min; (c) PM at 1100 °C for 15 min; and (d) SM at 1100 °C for 15 min.
cated. The first was processed by rolling at 850 °C from 16 to 5 mm in eight rolling passes; the second was rolled at the same temperature from 16 to 2 mm in four rolling passes. Due to the difference in the final thickness, the materials are hereafter referred to as PM (plate material) and SM (sheet material), respectively. The textures of the beta phase of the PM and SM materials, which would impact the subsequent beta grain growth due to the epitaxial growth of beta phase during the dissolution of alpha early in the annealing process, are shown in Fig. 6. The (110) pole figures for PM and SM were generally of the same type, but were much sharper in SM due to differences in the degree of continuous dynamic recrystallization (during rolling) and static discontinuous recrystallization (during interpass reheating) imparted by the thermomechanical processing of the two lots of material. Such textures, however, are generally typical of rolled bcc metals.
3. Experimental procedures Isothermal beta annealing of the PM and SM lots of Ti – 6Al –4V was conducted to determine the effect of texture on beta grain growth. To this end, samples of each material were heat treated in a vacuum furnace at temperatures of 1000, 1030, 1060, 1100, 1130, 1160, 1200, 1230, and 1260 °C for times ranging from 5 to 90 min. The relatively small increment in temperature (i.e. 30 °C) used in these experiments was chosen to provide sufficient data to ensure a very accurate determination of the grain-growth activation energy Q and graingrowth exponent n in Eq. (1) within the given temperature range. After isothermal exposure, each specimen was cooled rapidly by turning off the power to the water-cooled furnace. This yielded an initial sample cooling rate of not less than 50 K s − 1, which was sufficient to stop beta grain growth. At temperatures
O.M. I6asishin et al. / Materials Science and Engineering A332 (2002) 343–350
below the beta transus, the sample cooling rate was somewhat lower, thereby promoting diffusional decomposition of the beta phase. By this means, the beta grain boundaries were decorated with a thin layer of alpha that was beneficial in determining the beta grain size during subsequent optical metallography. Average grain sizes were determined by using the linear intercept method; grains close to the surface were not taken into account to avoid the pinning effect of the free surface. For both materials, the standard deviation was approximately 17–20% of the average grain size with a confidence interval (CI) of 95%.
4. Results and discussion Typical microstructures of the SM and PM materials after identical isothermal annealing treatments are compared in Fig. 7. In all cases, a uniform equiaxed beta grain structure was developed. Some quantitative grain-size measurements are summarized in Fig. 8; the average beta
Fig. 8. Average beta grain size as a function of annealing time at (a) 1000 °C; or (b) 1200 °C.
347
grain size as a function of annealing time at 1000 or 1200 °C is shown for the PM and SM materials. By and large, the maximum grain sizes were approximately 1.0–1.2 mm for PM and 0.3– 0.4 mm for SM, or values substantially smaller than the smallest cross section dimension (= the sheet per plate thickness) in each case. Grain size measurements for all of the annealing temperatures are shown in Fig. 9a (PM) and Fig. 9b (SM). An inspection of the data in Figs. 8 and 9 reveals that the PM material exhibited much faster beta grain growth than the SM material. Equally important, the experimental data for SM had marked indications of discontinuous beta grain growth characterized by a significant period of stagnation when the grain size was about 250– 300 mm. On the other hand, the PM material exhibited rather typical isothermal grain-growth behavior. Plots of ln D versus ln t (Figs. 10 and 11) were made for both materials to determine the value of the graingrowth exponent n and thus the basis for applying Eq. (1) to the experimental data. However, these graphs showed that n did not remain constant over the range of temperatures investigated in the present work. Moreover, the ln D versus ln t graphs were quite different for PM and SM. For the PM material (Fig. 10), grain growth was characterized by a constant value of n for each annealing temperature; the specific value, however, was temperature dependent (Fig. 12). A sigmoidal dependence n(T) was found, thus indicating slower grain growth at higher temperatures when the grain size was large. Despite the fact that the ln D versus ln(1/T) graph is linear for the PM material (Fig. 13), the evaluation of the activation energy Q could not be done because of the sigmoidal n(T) dependence. For the SM material, the n(T) dependence was even more complicated (Fig. 11). At low annealing temperatures, the slope of the ln D versus ln t plots (Fig. 11a) yielded an average n of approximately 8.9, or a value indicative of very slow grain growth. At temperatures above 1100 °C (Fig. 11b), the grain-growth rate increased with both temperature and time. The stage of grain-growth stagnation observed at these temperatures when the average beta grain size was about 250– 300 mm changed to relatively fast growth at longer times. The higher the temperature, the shorter was the exposure necessary for fast growth to start. Using data for 15–60 min heat treatments, an n value of 3.9 was derived for the higher-temperature behavior. With such an n value, the activation energy Q for SM was estimated as 227 kJ mol − 1 from the corresponding ln D versus ln(1/T) graph (Fig. 14). The observations from the present investigation show that materials with different initial textures exhibit noticeably different grain-growth kinetics during isothermal heat treatments. The data also revealed indications of discontinuous growth characterized by alternating stages of relatively fast (normal) grain growth and relatively slow growth (Fig. 9). This phenomenon was
348
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Fig. 9. Average beta grain size as a function of annealing time for (a) PM; or (b) SM.
Fig. 10. ln D versus ln t dependence for PM material at various temperatures.
apparent in SM at all temperatures. An initial period of fast growth during the first 5 min to a grain size of approximately 250 mm was followed by very slow growth and then a second stage of fast growth. The time interval of slow growth decreased as the annealing temperature
was raised. The annealing temperature also controlled the final grain size that was achieved by the end of a 1-h exposure; the final grain sizes were 400, 600, 750, and 900 mm at 1130, 1160, 1200, and 1230 °C, respectively. In the PM material, discontinuous grain growth was not as
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evident for the exposure times employed in the present work. An increase in grain growth rate at longer times was observed only at high temperatures; at these temperatures, the initial value of n was approximately 5.6, which itself corresponds to relatively slow growth. The present observations agree well with previous, less detailed measurements of discontinuous grain growth for
Fig. 13. ln D versus ln(1/T) for PM material at t= 15 min.
Fig. 14. ln D versus ln(1/T) for SM material at t = 15 min.
Fig. 11. ln D versus ln t dependence for SM material at various temperatures.
IMI685 and Ti–6Al –4V [3]. The experimental data can also be rationalized on the basis of previous analytical modeling [11] that suggests the possibility of discontinuous grain growth in initially strongly textured materials. This modeling work revealed that a relatively strong initial texture could be the most likely reason for very slow initial grain growth and much quicker growth after a reduction in the sharpness of the texture at long annealing times.
5. Conclusions
Fig. 12. Temperature dependence of the grain growth exponent n for PM material.
A series of carefully controlled, isothermal beta annealing trials were conducted on samples of Ti– 6Al – 4V processed to yield identical, fine-grained starting microstructures but different crystallographic textures. Measurements of beta grain growth kinetics led to the following conclusions: (1) Beta-grain growth under isothermal annealing conditions is strongly affected by texture in materials which are otherwise identical in
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terms of chemical composition and microstructure. (2) The kinetics of beta grain growth in titanium alloys can not be generally described in terms of a constant (temperature-independent) grain-growth exponent and activation energy. (3) Strict control of initial texture is required in thermomechanical processing of titanium alloys to ensure the desired level of those mechanical properties which are affected by the beta grain size.
Acknowledgements The present work was supported by the Air Force Office of Scientific Research (AFOSR) and the AFOSR European Office of Aerospace Research and Development (AFOSR/EOARD) within the framework of STCU Partner Project P-041. The encouragement of the AFOSR program managers (Dr R.S. Fredell and Dr C.S. Hartley) is greatly appreciated.
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