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Effect of heat treatment and hot isostatic pressing on void density and fracture mode of Al67Ni8Ti25
Scripta METALLURGICA
Vol. 23, pp. 1827-1830, 1989 Printed in the U.S.A.
Pergamon Press plc All rights reserved
EFFECT OF HEAT TREATMENTAND HOT ISOSTATIC PRESSINGON VOID DENSITY AND FRACTUREMODEOF Al67Ni8Ti25 D.D. Mysko*, J.B Lumsden+, W.O. Powers$ and J.A. Weft* * Department of Materials Science University of Virginia, Charlottesville, VA 22901 + Rockwell International Science Center 1049 Camino dos Rios, Thousand Oaks, CA 91360 $ Now at Olin Corp., Metals Research Laboratory 91Shelton Ave., New Haven, CT 06511 (Received July 28, 1989)
Introduction Raman and Schubert [1] f i r s t reported that the tetragonal D022 crystal structure of Al~Ti can be changed to the related cubic LI 2 crystal structure by substi[ution of Ni for Al to achieve compositions near A167NiATi~. - The same crystal structure change has been shown to occur for substitution of Cu;-Zn; F~T Pd, and Co for part of the Al in Al~Ti [I-4]. Following the suggestion of Nash et al [5], the LI 2 phases in these ternary alloy s~stems will be denoted phase throughout the present paper. Several investigators have examined the deformation and fracture behavior of ~ phase alloys [6-10]. All published reports of the fracture characteristics of x phase alloys agree that these alloys fail in an extremely b r i t t l e fashion by transgranular cleavage. In A167NIRTi25, cleavage has been reported to occur on a variety of crystallographic planes, including-{100}, {110}, {111} and others [9,10]. Only one measurement of toughness has been published; Kic = 3 MPaJm for A167NiBTi25 at room temperature [10]. Homogenization heat treatment of conventionally-solidified x phase alloys is necessary to reduce the volume fraction of nonequilibrium second phases. Voids have been observed in the microstructures of many of the x phase alloys that have been homogenized after s o l i d i f i c a t i o n [4,9,10]. The origin of the voids and their possible participation in b r i t t l e fracture processes have been the subjects of speculation by previous investigators [10], but few conclusive experimental results have been reported. The purpose of the present paper is to describe and to interpret experimental observations pertaining to void formation and to the effect of voids on fracture in A167Ni8Ti25. Procedure Button-shaped samples with nominal composition A167NiRTip~ were prepared from pure elements by non-consumable electrode arc-melting. SpecTmenEc~[ from the buttons by electrodischarge machining were encapsulated in s i l i c a tubes backfilled with argon gas. Homogenization heat treatment of the encapsulated specimens was carried-out by heating to 1323 K at a rate of 100 K/h, holding for lO0 h at 1323 K, then cooling to ambient temperature at a rate of 100 K/h. Hot Isostatic Pressing (HIPing) treatments were carried-out by heating previously-homogenized specimens to 1323 K at a rate of lO00 K/h, holding at 1323 K'under pressure for the desired period, and cooling the specimens to ambient temperature at a rate of approximately lO00 K/h. Point counting methods were used to measure the void and second phase volume fractions. A specimen of A167Ni8Ti25 was fractured in an ultrahigh vacuum chamber
1827 0036-9748/89 $3.00 + .00 Copyright (c) 1989 Pergamon Press plc
1828
FRACTURE MODE OF AI-Ni-Ti
Vol.
25, No. Ii
TABLE 1. Second Phase and Void Volume Fractions in Al67Ni8Ti25
Condition and Treatment
Second phase volume fraction*
As-cast Homogenized 1323 K / 100 h HIP 1323 K / 1 h / 140 MPa HIP 1323 K / 18 h / 200 MPa
0.10 ± 0.045 ± 0.038 ± 0.041 ±
0.023 0.015 0.014 0.016
Void volume fraction* None observed 0.060 ± 0.017 0.055 ± 0.017 0.009 ± 0.007
Confidence intervals represent lo for a normal distribution to detect possible hydrogen gas emission from the voids. The hydrogen pressure in the chamber was monitored using a quadrapole mass spectrometer precision gas analyzer (model UTI IOOC). Observations Figure I shows the microstructure of the AI67NiRTi)B alloy after solidification and after homogenization heat treatment. Thesemicrographs%ho~ tEAt the void volume fraction is increased substantially by the homogenization treatment. Examination of many cross-sections of the homogenized A167NisTi2s alloy revealed that the void volume fraction is variable; the micrographs shown ih FiguP~ I are chosen to represent typical void distributions in the alloy. Figure I also demonstrates that HIP±rig previously-homogenized AI67NiRTi2s can appreciably reduce the void volume fraction. Table I lists the second phase-and-void volume fractions in the as-cast condition, after homogenization, and after various HIPing treatments, which were measured from representative micrographs such as those shown in Figure I. In the homogenized condition, the average void diameter on the plane of polish is 8.6 /an. Assuming the vo)~s a~e monodispersed spheres, the true void diameter is 14 pm and the void density is 4.4 x 10*° m"*. The effect of void volume fraction on the fracture surface morphology of bend specimens has been determined. Figure 2 shows that reducing the void volume fraction has no discernible effect on the fracture mode; fracture occurs by transgranular cleavage in both cases. This fracture mode is associated with very low toughness in homogenized A167Ni8Ti25 [10].
i
Figure I. Mlcrostructure of AI~7NiRTI)B after a variety of treatments. (a) As-cast. (b) After homogenization at 1323 K for I00 hT Ic) After HIPing at 1323 K and 200 MPa for 18 h.
Vol.
23, No.
Ii
FRACTURE MODE OF A I - N i - T i
1829
~ }
Figure 2. Fracture morphology of AI67Ni8Ti25. at ]323 K and 200 MPa for 18 h.
(a) Homogenized at 1323 K for 100 h.
(b) HIPed
A mechanism previously suggested for the substantial increase in void volume fraction during homogenization is precipitation of hydrogen gas incorporated into the alloy during melting [10]. (In this case, the voids would be more correctly termed "pores"). In an effort to detect the presence of hydrogen gas in the voids, a homogenized sample of A167NisTi25 was Fractured in an ultrahigh vacuum chamber while monitoring the hydrogen pressure-~n the-Ehamber. The total residual gas pressure was 6 x 10-° Pa when the sample was fractured; no change in pressure of hydrogen gas was detected upon fracture of the Al67Ni8Ti25 sample. The s e n s i t i v i t y of the gas analyzer was sufficient to detect release of 3 x_lO-13 moles of hydrogen gas. The cross-sectional area of the fractured sample was 2 x 10-b mZ. Based on void volume fraction and siI~ m~asurements, the void volume e~posed by fracturing the sample was estimated to be 2 x 10-~ m~. Using an estimate of 2 J/m~ for the free surface energy at the homogenization temperature [10], the equilibrium gas pressure ~n the voids should be 600 kPa and the amount of gas released upon fracture should be 5 x 10-~ moles. These estimates show that the s e n s i t i v i t y of the gas analyzer was more than sufficient to detect the amount of hydrogen gas that would be required to form the observed void population by precipitation from solid solution during annealing. Discussion
Three potential mechanisms for void formation in A167NisTi25 have been discussed by previous investigators [4,9,10]: precipitation of gaseous ~mp~rities incorporated into the alloy during melting, a volume change associated with elimination of second phases during homogenization, and Kirkendall voids associated with diffusion during homogenization. The operative mechanism can be inferred from the experimental results of the present investigation. i . No hydrogen gas was detected escaping from the voids when they were exposed by fracture. We conclude that the voids are not formed by precipitation of hydrogen during annealing. i i . The increase in void volume fraction upon homogenization is approximately equal to the decrease in second phase volume fraction. I f a volume change associated with elimination of second phase particles was responsible for void formation, the entire volume of second phase particles eliminated would have to be converted to voids to account for the observed volume fraction changes. This is unreasonable. i i i . No direct results in support of the Kirkendall mechanism of void formation have been obtained for Al~TNi~Ti~. However, recent studies of TiAI-Nb diffusion couples have shown substantial d]ff~Fences in Al and Ti mobilities [11]. I t is not unreasonable to suppose that similar differences could exist in Al~7Ni~Tips. The above arguments lead to the conclusion that void forEkti6n 6~ a Kirkendall mechanism during homogenization of A167NiBTi25 is the only proposed mechanism that is consistent with experimental results.
1830
FRACTURE MODE OF AI-Ni-Ti
Vol.
23, No. ii
Observations reported in the present paper show that the fracture mode is unchanged by substantially reducing the void volume fraction by HIPing. In a previous paper, Turner et al [10] concluded that cleavage fracture of A167Ni8Ti25 is the result of an activation barrier for dislocation emission from crack tips, which prevenls plastic blunting of the cracks [12]. The present results are consistent with the explanation proposed by Turner et al [10], since microstructural heterogeneities on the scale of the void dispersion observed in Al67Ni8Ti25 would not be expected to affect dislocation emission from crack tips. Acknowledqements Experimental assistance of B. Myers is gratefully acknowledged. This work was supported by the Wright Aeronautical Laboratories under contract number F3361S-86-C-SO74, S.D. Kirchoff was the technical monitor. References I. 2. 3. 4.
A. Raman and K. Schubert, Z. Metallkde., 56 (1965) 99. V.Ya. Markiv and V.V. Burnashova, Metallofis. Akad. Nauk. Ukrainokoj SSR, 46 (1973) 103. A. Siebold, Z. Metallkde., 72 (1981) 712. W.O. Powers and J.A. Wert, "Microstructure, Deformation and Fracture Characteristics of an AI~7PdRTi~ Intermetallic Alloy", accepted for publication in Met. Trans. 5. P.{~ N~shT-V. Vejins and W.W. Liang, Bulletin of Alloy Phase Diagrams, 3 (1982) 367. 6. S.C. Huang, E.L. Hall and M.F.X. Gigliotti, J. Mater. Res., 3 (1988) I. 7. K.S. Kumar and J.R. Pickens, in "Dispersion Strengthened Aluminum Alloys", Y.-W. Kim and W. Griffith (eds.), TMS, Warrendale, 1988, pp. 763-786. 8. K.S. Kumar and J.R. Pickens, Scripta Met., 22 (1988) 1015. 9. E.P. George, W.D. Porter, H.M. Henson and W.C. Oliver, J. Mater. Res., 4 (1989) 78. 10. C.D. Turner, W.O. Powers and J.A. Wert, "Microstructure, Deformation and Fracture Characteristics of an A167NiBTi25 Intermetallic Alloy", Acta Met., in press. 11. T.J. Jewett, J.C. Lin, N~R. Bonda, L.E. Seitzman, K.C. Hsieh, Y.A. Chang and J.H. Perepezko, in "High Temperature Ordered Intermetallic Alloys", C.C. Koch, C.T. Liu, N.S. Stoloff and A.I. Taub (eds), MRS, in press. 12. J.R. Rice and R. Thomson, Phil. Mag., 29 (1974) 73.
Vol. 23, pp. 1827-1830, 1989 Printed in the U.S.A.
Pergamon Press plc All rights reserved
EFFECT OF HEAT TREATMENTAND HOT ISOSTATIC PRESSINGON VOID DENSITY AND FRACTUREMODEOF Al67Ni8Ti25 D.D. Mysko*, J.B Lumsden+, W.O. Powers$ and J.A. Weft* * Department of Materials Science University of Virginia, Charlottesville, VA 22901 + Rockwell International Science Center 1049 Camino dos Rios, Thousand Oaks, CA 91360 $ Now at Olin Corp., Metals Research Laboratory 91Shelton Ave., New Haven, CT 06511 (Received July 28, 1989)
Introduction Raman and Schubert [1] f i r s t reported that the tetragonal D022 crystal structure of Al~Ti can be changed to the related cubic LI 2 crystal structure by substi[ution of Ni for Al to achieve compositions near A167NiATi~. - The same crystal structure change has been shown to occur for substitution of Cu;-Zn; F~T Pd, and Co for part of the Al in Al~Ti [I-4]. Following the suggestion of Nash et al [5], the LI 2 phases in these ternary alloy s~stems will be denoted phase throughout the present paper. Several investigators have examined the deformation and fracture behavior of ~ phase alloys [6-10]. All published reports of the fracture characteristics of x phase alloys agree that these alloys fail in an extremely b r i t t l e fashion by transgranular cleavage. In A167NIRTi25, cleavage has been reported to occur on a variety of crystallographic planes, including-{100}, {110}, {111} and others [9,10]. Only one measurement of toughness has been published; Kic = 3 MPaJm for A167NiBTi25 at room temperature [10]. Homogenization heat treatment of conventionally-solidified x phase alloys is necessary to reduce the volume fraction of nonequilibrium second phases. Voids have been observed in the microstructures of many of the x phase alloys that have been homogenized after s o l i d i f i c a t i o n [4,9,10]. The origin of the voids and their possible participation in b r i t t l e fracture processes have been the subjects of speculation by previous investigators [10], but few conclusive experimental results have been reported. The purpose of the present paper is to describe and to interpret experimental observations pertaining to void formation and to the effect of voids on fracture in A167Ni8Ti25. Procedure Button-shaped samples with nominal composition A167NiRTip~ were prepared from pure elements by non-consumable electrode arc-melting. SpecTmenEc~[ from the buttons by electrodischarge machining were encapsulated in s i l i c a tubes backfilled with argon gas. Homogenization heat treatment of the encapsulated specimens was carried-out by heating to 1323 K at a rate of 100 K/h, holding for lO0 h at 1323 K, then cooling to ambient temperature at a rate of 100 K/h. Hot Isostatic Pressing (HIPing) treatments were carried-out by heating previously-homogenized specimens to 1323 K at a rate of lO00 K/h, holding at 1323 K'under pressure for the desired period, and cooling the specimens to ambient temperature at a rate of approximately lO00 K/h. Point counting methods were used to measure the void and second phase volume fractions. A specimen of A167Ni8Ti25 was fractured in an ultrahigh vacuum chamber
1827 0036-9748/89 $3.00 + .00 Copyright (c) 1989 Pergamon Press plc
1828
FRACTURE MODE OF AI-Ni-Ti
Vol.
25, No. Ii
TABLE 1. Second Phase and Void Volume Fractions in Al67Ni8Ti25
Condition and Treatment
Second phase volume fraction*
As-cast Homogenized 1323 K / 100 h HIP 1323 K / 1 h / 140 MPa HIP 1323 K / 18 h / 200 MPa
0.10 ± 0.045 ± 0.038 ± 0.041 ±
0.023 0.015 0.014 0.016
Void volume fraction* None observed 0.060 ± 0.017 0.055 ± 0.017 0.009 ± 0.007
Confidence intervals represent lo for a normal distribution to detect possible hydrogen gas emission from the voids. The hydrogen pressure in the chamber was monitored using a quadrapole mass spectrometer precision gas analyzer (model UTI IOOC). Observations Figure I shows the microstructure of the AI67NiRTi)B alloy after solidification and after homogenization heat treatment. Thesemicrographs%ho~ tEAt the void volume fraction is increased substantially by the homogenization treatment. Examination of many cross-sections of the homogenized A167NisTi2s alloy revealed that the void volume fraction is variable; the micrographs shown ih FiguP~ I are chosen to represent typical void distributions in the alloy. Figure I also demonstrates that HIP±rig previously-homogenized AI67NiRTi2s can appreciably reduce the void volume fraction. Table I lists the second phase-and-void volume fractions in the as-cast condition, after homogenization, and after various HIPing treatments, which were measured from representative micrographs such as those shown in Figure I. In the homogenized condition, the average void diameter on the plane of polish is 8.6 /an. Assuming the vo)~s a~e monodispersed spheres, the true void diameter is 14 pm and the void density is 4.4 x 10*° m"*. The effect of void volume fraction on the fracture surface morphology of bend specimens has been determined. Figure 2 shows that reducing the void volume fraction has no discernible effect on the fracture mode; fracture occurs by transgranular cleavage in both cases. This fracture mode is associated with very low toughness in homogenized A167Ni8Ti25 [10].
i
Figure I. Mlcrostructure of AI~7NiRTI)B after a variety of treatments. (a) As-cast. (b) After homogenization at 1323 K for I00 hT Ic) After HIPing at 1323 K and 200 MPa for 18 h.
Vol.
23, No.
Ii
FRACTURE MODE OF A I - N i - T i
1829
~ }
Figure 2. Fracture morphology of AI67Ni8Ti25. at ]323 K and 200 MPa for 18 h.
(a) Homogenized at 1323 K for 100 h.
(b) HIPed
A mechanism previously suggested for the substantial increase in void volume fraction during homogenization is precipitation of hydrogen gas incorporated into the alloy during melting [10]. (In this case, the voids would be more correctly termed "pores"). In an effort to detect the presence of hydrogen gas in the voids, a homogenized sample of A167NisTi25 was Fractured in an ultrahigh vacuum chamber while monitoring the hydrogen pressure-~n the-Ehamber. The total residual gas pressure was 6 x 10-° Pa when the sample was fractured; no change in pressure of hydrogen gas was detected upon fracture of the Al67Ni8Ti25 sample. The s e n s i t i v i t y of the gas analyzer was sufficient to detect release of 3 x_lO-13 moles of hydrogen gas. The cross-sectional area of the fractured sample was 2 x 10-b mZ. Based on void volume fraction and siI~ m~asurements, the void volume e~posed by fracturing the sample was estimated to be 2 x 10-~ m~. Using an estimate of 2 J/m~ for the free surface energy at the homogenization temperature [10], the equilibrium gas pressure ~n the voids should be 600 kPa and the amount of gas released upon fracture should be 5 x 10-~ moles. These estimates show that the s e n s i t i v i t y of the gas analyzer was more than sufficient to detect the amount of hydrogen gas that would be required to form the observed void population by precipitation from solid solution during annealing. Discussion
Three potential mechanisms for void formation in A167NisTi25 have been discussed by previous investigators [4,9,10]: precipitation of gaseous ~mp~rities incorporated into the alloy during melting, a volume change associated with elimination of second phases during homogenization, and Kirkendall voids associated with diffusion during homogenization. The operative mechanism can be inferred from the experimental results of the present investigation. i . No hydrogen gas was detected escaping from the voids when they were exposed by fracture. We conclude that the voids are not formed by precipitation of hydrogen during annealing. i i . The increase in void volume fraction upon homogenization is approximately equal to the decrease in second phase volume fraction. I f a volume change associated with elimination of second phase particles was responsible for void formation, the entire volume of second phase particles eliminated would have to be converted to voids to account for the observed volume fraction changes. This is unreasonable. i i i . No direct results in support of the Kirkendall mechanism of void formation have been obtained for Al~TNi~Ti~. However, recent studies of TiAI-Nb diffusion couples have shown substantial d]ff~Fences in Al and Ti mobilities [11]. I t is not unreasonable to suppose that similar differences could exist in Al~7Ni~Tips. The above arguments lead to the conclusion that void forEkti6n 6~ a Kirkendall mechanism during homogenization of A167NiBTi25 is the only proposed mechanism that is consistent with experimental results.
1830
FRACTURE MODE OF AI-Ni-Ti
Vol.
23, No. ii
Observations reported in the present paper show that the fracture mode is unchanged by substantially reducing the void volume fraction by HIPing. In a previous paper, Turner et al [10] concluded that cleavage fracture of A167Ni8Ti25 is the result of an activation barrier for dislocation emission from crack tips, which prevenls plastic blunting of the cracks [12]. The present results are consistent with the explanation proposed by Turner et al [10], since microstructural heterogeneities on the scale of the void dispersion observed in Al67Ni8Ti25 would not be expected to affect dislocation emission from crack tips. Acknowledqements Experimental assistance of B. Myers is gratefully acknowledged. This work was supported by the Wright Aeronautical Laboratories under contract number F3361S-86-C-SO74, S.D. Kirchoff was the technical monitor. References I. 2. 3. 4.
A. Raman and K. Schubert, Z. Metallkde., 56 (1965) 99. V.Ya. Markiv and V.V. Burnashova, Metallofis. Akad. Nauk. Ukrainokoj SSR, 46 (1973) 103. A. Siebold, Z. Metallkde., 72 (1981) 712. W.O. Powers and J.A. Wert, "Microstructure, Deformation and Fracture Characteristics of an AI~7PdRTi~ Intermetallic Alloy", accepted for publication in Met. Trans. 5. P.{~ N~shT-V. Vejins and W.W. Liang, Bulletin of Alloy Phase Diagrams, 3 (1982) 367. 6. S.C. Huang, E.L. Hall and M.F.X. Gigliotti, J. Mater. Res., 3 (1988) I. 7. K.S. Kumar and J.R. Pickens, in "Dispersion Strengthened Aluminum Alloys", Y.-W. Kim and W. Griffith (eds.), TMS, Warrendale, 1988, pp. 763-786. 8. K.S. Kumar and J.R. Pickens, Scripta Met., 22 (1988) 1015. 9. E.P. George, W.D. Porter, H.M. Henson and W.C. Oliver, J. Mater. Res., 4 (1989) 78. 10. C.D. Turner, W.O. Powers and J.A. Wert, "Microstructure, Deformation and Fracture Characteristics of an A167NiBTi25 Intermetallic Alloy", Acta Met., in press. 11. T.J. Jewett, J.C. Lin, N~R. Bonda, L.E. Seitzman, K.C. Hsieh, Y.A. Chang and J.H. Perepezko, in "High Temperature Ordered Intermetallic Alloys", C.C. Koch, C.T. Liu, N.S. Stoloff and A.I. Taub (eds), MRS, in press. 12. J.R. Rice and R. Thomson, Phil. Mag., 29 (1974) 73.