- Email: [email protected]
Densification and microstructural control of near gamma;-TiAl intermetallic powders by HIPing
HOT ISOSTATIC PRESSING ’93 L. Delaey and H. Tas (Editors) 1994 Elsevier Science B.V.
99
Densification and microstructural control of near γ-TiAl intermetallic powders by HIPing W. Wallace*, L. Zhao* J.C. Beddoesk and D. Morphy* aStructures and Materials Laboratory, Institute for Aerospace Research, National Research Council of Canada, Ottawa, Ontario, K1A 0R6 ^Department of Mechanical and Aerospace Engineering, University, Ottawa, Ontario, K1A 5B6
Carleton
A bstract
Hot isostatic pressing has been used to consolidate pre-alloyed powders of near γ-titanium aluminide with attention being paid to the effects of HIP parameters on densification, microstructural control and properties. Three alloys have been studied, namely Ti-48A1, Ti-48A1-2W, and Ti-47.5Al-3Cr. (atom %). The W containing alloy required a temperature of 1100°C to eliminate porosity, but full density was achieved at lower temperatures for the other two alloys. While the consolidated Cr powder possessed a uniform microstructure, the other two alloys retained a dendritic structure even when HIPed at 1250°C. HIP cycles with varying temperatures and pressures were therefore designed to refine and homogenize the microstructures. The densities and microstructures resulting from the modified HIP cycles were characterized and compared to those obtained at constant temperatures and pressures. Methods of optimizing the HIP parameters to produce a homogenous microstructure suitable for further thermo-mechanical processing (TMP) or service application are discussed. 1. INTRODUCTION
The popular appeal of hot isostatic pressing (HIP) lies in its ability to heal porosity or similar defects in metal or ceramic parts or preforms. However the full potential of HIP lies in its ability to produce parts having microstructures tailored to meet the design requirements of parts in an optimum manner. This has been demonstrated at NRC and elsewhere for superalloy powder compacts and castings (1-3). Similar work, directed at the optimum HIP processing of TiAl is being carried out at NRC (4-6), and by other groups such as that of Schaeffer (7-9). In this paper the effects of HIP on the microstructure and properties of three near γ-TiAl powders are examined with particular attention to the mechanisms of densification and suitability of the compacts for their intended applications.
100
2. EXPERIMENTAL DETAILS
The powders were prepared by Crucible Compaction Metals Inc. (Oakdale, Pennsylvania) using the titanium gas atomization process (TGA), (10). The actual chemical compositions of the binary and ternary powders are Ti-48.1A1, Ti-48.4Al-1.9W and Ti-47.1Al-2.9Cr, respectively. The powders were HIP’ed in evacuated, out-gassed and sealed stainless steel or titanium containers (4). Isothermal hot compression tests were carried out on the ternary alloys in the temperature range 900-1200°C and at constant true strain rates in the range 10 4sec-i and lO -isec1. Ambient and elevated temperature tensile tests were also carried out to determine flow and fracture properties. Three types of HIP cycle were employed in the study: a) Cycle type 1: Ramping to constant temperature and pressure periods of 120 minutes, followed by in-autoclave cooling. Maximum temperature was in the range 1050°C and 1250°C, and peak pressure was 200 MPa. b) Cycle type 2: Ramping to 1380°C/200 MPa/30 minutes followed by 1100°C/170 MPa/90 minutes and in-autoclave cooling. c) Cycle type 3: Ramping to 1100°C/200 MPa/110 minutes followed by ramping to peak temperatures of 1400°C and 1415°C/200 MPa/10 minutes. Treatments (b) and (c) have only been performed on the W-containing alloy. All three treatments involved a total time of 2 hours at the hold temperatures. 3. RESULTS AND DISCUSSION 3 .1 A s-r ece iv ed p ow d ers
The binary and the Ti-48A1-2W powders had a well defined dendritic structure, whereas the Cr-containing powder was far more uniform in microstructure. The two ternary powders are compared in Figure 1, where it can be seen that light etching dendrite cores are clearly visible and quite abundant in the W-alloy. These cores are W-rich, hard and extremely stable. They resist deformation and cause plastic flow during HIP or subsequent tensile or compressive straining to be concentrated in the softer interdendritic regions. This coring is likely to inhibit plasticity and hence densification during HIP, and restrict ductility and superplasticity during subsequent tensile deformation or forging of the compacts. The high temperature treatments of HIP cycles 1 and 2 were attempts to eliminate this dendritic segregation and improve flow and fracture properties. The principal phase in the Ti-48A1 powder was γ -TiAl with a small amount (-15%) of a2. The principal phase in the W and Cr containing powders was 0C2 , with about 40% γ and about 10%βο, as illustrated for the W-containing alloy in Figure 2.
101
Figure 1. Microstructures of as-received powders, a) Ti-48A1-2W and b) Ti-47.5A1-Cr. 0.50 021 a2
As-Received Ti-48A4-2W Powder
CO
b 111-v
,
0020C 2
m 025
CL
O
,
020a2
10b
02 2 ,
L 35
40
" 7
002ώ 00γ 45 2Θ
50
55
Figure 2. X-ray diffraction plot of Ti-48A1-2W as-received powder. 3 .2 HIP c y c le s ty p e 1
The HIP temperatures were such that all alloys should be in an (X2 +J (1050 or 1100°C) or α+γ (1150° or 1250°C) condition. Full density (>99.9%) was achieved in the binary alloy and the Cr-alloy at the lowest HIP temperature employed, whereas a temperature of 1100°C was needed for the W-containing alloy. The microstructures of these two ternary materials after HIPing at 1050°C/200MPa are shown in Figure 3. Prior particle boundaries (PPB’s) can be seen in both compacts, but while the W-alloy shows an unrecrystallized structure with strong dendritic segregation, the Cr-containing alloy shows extensive recrystallization with profuse twinning. PPB’s in these alloys tend to provide paths for easy crack propagation, and hence they may impede ductility. The higher HIP temperatures, up to 1250°C can partially eliminate the PPB’s and improve homogeneity and
102
Figure 3. As-HIP 1050°C/200MPa, a) Ti-48A1-2W, b) Ti-47.5Al-3Cr.
A
Figure 4. As-HIP 1250°C/200MPa, a) Ti-48A1-2W, b) Ti-47.5Al-3Cr. ductility. However, they do not completely eliminate dendritic segregation in the W-containing alloy. The best combination of room temperature strength and ductility in the as-pressed W-alloy was 505MPa and 1.3% after 1250°C HIP (5). The microstructures of the alloys after HIPing at 1250°C/200MPa are shown in Figure 4. This higher temperature promotes recrystallization and increases grain size, as shown for the W-alloy in Table 1. In terms of phase content, both HIP temperatures of 1050°C and 1250°C cause the transformation of 0C2 to γ, as illustrated for the W-containing alloy by comparing Figure 5 with Figure 2. In this alloy after 1250°C HIP, the ßQ phase is found predominantly in the dendrite cores as small recrystallized grains, Figure 6a. In the Crcontaining alloy after 1250°C HIP, the β0 phase is fairly uniformly distributed throughout the compact, and is located at triple points and grain boundaries of the γ-phase, Figure 6b.
103
2.50
CO
ο
* 1.25
CO Q_
O
35
40
45 2Θ
50
55
Figure 5. X-ray diffraction plot of HIPped Ti-48A1-2W compact.
Table 1 Summary of the response of Ti-48A1-2W powder to Hot Isostatic Pressing No.
1. 2. 3. 4.
HIP Cycle
200MPa/1050°C/2hr 200MPa/1100°C/2hr 200MPa/1150°C/2hr 200MPa/1250°C/2hr
5. 200MPa/1380°C/0.5hr +170MPa/1100°C/1. 5hr 6. 200MPa/1100°C/ Ihr 50 min +200 MPa/ 1400°C/10 min 7. 200MPa/1100°C/lhr 50 min +200MPa/1415°C /10 min
Prior Dendritic Structure
Microstructural Average Grain Densification Size (μ) Homogeneity
retained retained retained retained
poor poor poor improved
4.1 6.0 8.0 12.5
>99.5% >99.9% >99.9% >99.9%
removed
improved
--
>99.9%
removed
removed
~200μ/Ι>25μ/γ Good (L+γ) L=ot2 +Y lam ellar Good (near fully L)
~300μ/Ι>
>99.9%
>99.9%
While mechanical properties are improved by HIPing at higher temperatures, it is unlikely that these compacts would be used for engineering components because of their low ductility. Isothermal forging is being explored as a means of refining microstructures and achieving a more favourable combination of properties. Forging preforms are therefore required that have low flow stress, and good ductility. In this case the use of high HIP temperatures is deleterious since it reduces the range of temperatures and strain rates over which superplasticity is observed, as shown in Figure 7. In this figure the contours show lines of constant strain rate sensitivity exponent m, obtained from the slope of
104
Figure 6. SEM images of ß 0 phase in a) Ti-48A1-2W and b) Ti-47.5Al-3Cr as-HIP 1250°C/200MPa. plots of log stress (σ) versus log strain rate (έ) at constant strain and temperature, according to the relationship σ = dm 3 .3 D en sifica tio n
The rates of densification of TiAl powders have been studied experimentally by Schaeffer et al. (9) using an eddy current sensor which surrounds the shrinking container. The present results are qualitatively similar to those of Schaeffer and are consistent with the models of Helle et al. (11). Initial densification occurs rapidly by plastic flow due to stress
TEST TEMPERATURE (°C)
TEST TEMPERATURE (<>C)
Figure 7. Plots of constant strain rate sensitivity exponent for Ti-48A1-2W HIPped at a)1050°C/200MPa and b) 1250°C/200MPa.
105
I
F ig u re 8 . A s-H IP ped Ti-48A1-2W , 1 1 0 0 ° C /2 0 0 M Pa. c o n c e n tr a tio n s a t p o in ts of c o n ta c t b e tw e e n p a r tic le s , a s d e s c rib e d b y H e w itt e t al. (12). E v id e n c e of p la s tic flow is fo u n d a t th e lo w e r H IP te m p e r a tu r e s in th e fo rm of d e fo rm a tio n tw in s, a n d d is lo c a tio n s e g m e n ts a n d ta n g le s a s sh o w n in Fig. 8 . T h e se fe a tu re s a re fo u n d p re d o m in a n tly in th e s o ft in te r d e n d r itic re g io n s of th e W -alloy w h e re a s th e h a r d d e n d r ite co res re m a in re la tiv e ly u n d e fo rm e d . A s th e H IP te m p e r a tu r e is in c re a s e d to 12 5 0 °C , d y n a m ic r e c r y s ta lliz a tio n o c c u rs th r o u g h o u t th e c o m p a c ts in c lu d in g th e d e n d rite co res, a n d th is e lim in a te s m o s t of th e d is lo c a tio n s . T ab le 2 E s tim a te s o f th e d e n s ity a c h ie v e d b y p la s tic y ie ld in g in T i-48A 1-2W p o w d e r a s a fu n c tio n o f HIP te m p e r a tu re In itia l D e n sity (Do)
H IP C o n d itio n s °C /M P a(P )
F low S tr e s s (ay), M P a
Yield in d u c e d d e n s ity (Dy)
0 .6 4 0 .6 4 0 .6 4
1 0 5 0 /2 0 0 1 2 5 0 /2 0 0 1 4 0 0 /2 0 0
350 150 80
0.86
0 .7 5 0 .9 8 5
C a lc u la te d from H elle e t al (11).
T h e e q u a tio n s of H elle e t al. (11) c a n b e u s e d to c a lc u la te th e in itia l d e n s if ic a tio n t h a t o c c u r s b y p la s tic y ie ld in g , a n d t h i s p ro v id e s a n in d ic a tio n o f th e r e s id u a l p o ro s ity t h a t m u s t b e e lim in a te d b y c re e p o r o th e r m e c h a n is m s . T h e r e s u lts of th e s e c a lc u la tio n s a re s u m m a riz e d in
106
Table 2, where it can be seen that even a temperature as high as 1400°C would not achieve a density of 99% by plastic flow alone. In these calculations the initial density of the compacts was taken to be 0.64, the initial strain rate was estimated from the results of Schaeffer to be about 2.7 x 10-2 s-i and values of flow stress were obtained from the true stress versus true strain compression tests at the appropriate temperature and strain rate The residual porosity to be closed by creep or diffusion therefore varies from about 0.25 at 1050°C to about 0.015 at 1400°C, and evidence of creep is seen in Figure 8 in the form of dislocation sub-boundaries (marked “S”). A further mechanism of densification to consider is that of transformation induced plasticity, as proposed by Schaeffer and Janowski (8). They have suggested that in powders consisting predominantly of 2 phase, the transformation to γ will result in the formation of metastable 0C stacking faults and the passage of dislocations through the 0C2 phase, which contribute to densification by local deformation in the region of powder particle contact. The crystallographic relationships maintained by this transformation are those determined by Blackburn (13), such that the two phases share close packed planes and directions, (0002)α2//(111)γ and <1120>α2//<110>γ. As illustrated by comparing Figures 2 and 5, a substantial transformation of 0C2 to γ occurs in the W-containing powder, and therefore transformation induced plasticity as described by Schaeffer and Janowski is a possible contributor to densification in this alloy. 3 .4 HIP c y c le ty p e 2
The purpose of this treatment was to eliminate dendritic segregation while allowing only limited grain growth. The resulting microstructure contained large areas of equiaxed and fine-grained phases owing to
Figure 9. Ti-48A1-2W a) cycle 2, b) cycle 3 to 1400°C, c) cycle 3 to 1415°C.
107
complete recrystallization. Moreover, the dendritic structure no longer existed indicating that segregation was unstable under this HIP condition. While the overall microstructural homogenity was improved, regions containing coarse grains of γ-phase existed, Fig. 9a (arrows), with a typical size close to that of prior particles. This was attributed to the strong dendritic segregation in some powder particles where the interdendritic region dominated by γ single phase coarsened rapidly during HIPing due to the lack of second phase (4). Mechanical properties have not been determined for this structure. 3 .5 HIP c y c le ty p e 3
In these cycles the two stage process of cycle 2 was reversed to achieve initial densification and homogenization in the α2+γ phase field at 1100°C before completing the transformation to a at 1400°C. A thermocouple was mounted on the HIP can, and the can was packed in sand to ensure a slow heat up to the peak temperature. The microstructure, Figure 9b, consisted of a mixture of fine, equiaxed grained γ, and coarser grained lamellar α2+γ structure of the type desired. For this composition (Ti-48A12W), 1400°C is just above the a transus (14), but it appears that a fully transformed a structure was not achieved. In the latest run, the peak temperature was raised to 1415°C but the can was not packed in sand. Figure 9c shows that an almost fully transformed lamellar structure of α2+γ was obtained (>90%) but a small amount of equiaxed γ was retained. It is possible that the material did not experience the desired temperature of 1415°C for sufficiently long time to complete the transformation. 4.
FUTURE WORK
This work is continuing to produce a fully transformed α2+γ structure at which time the mechanical properties of the as-HIP material will be determined. Work is also continuing on the isothermal forging response of the Ti-48Al-3Cr alloy, which will determine the envelope for superplastic forging and elucidate the role of the β0 phase, Fig. 6b. 5.
CONCLUSIONS
• The presence of W in TiAl powders promotes dendritic segregation which is stable and impedes densification by plastic yielding. Chromium additions provide greater homogeneity and enhanced densification. • Densification occurs by plastic yielding, creep and transformation induced plasticity (α2->γ) in the W alloy, and probably the Cr alloy.
108
• High HIP temperatures in the α+γ phase field achieve full density and improved strength and ductility, but they reduce the envelope for superplastic flow during forging. • Even higher HIP temperatures in the single phase a field are feasible and should allow the formation of fully transformed α2+γ lamellar structures. 6 . ACKNOWLEDGEMENTS
This work was supported by the National Research Council of Canada under IAR-SML project JHR-01. Financial assistance was also provided by the Department of National Defence via Financial Arrangement 220791NRC51. The support of both organizations is gratefully acknowledged. 7 . REFERENCES
1 W. Wallace, J-P. Immarigeon and J.M. Trenouth, AGARD C.P. 200, Neuilly-sur-Seine (1976). 2 W. Wallace, H.B. Dunthome and R. Sprague, Can. Met. Quart., 13, No. 3 (1974), 517. 3 J.C. Beddoes and W. Wallace, Metallography, 13 (1980) 185. 4 J.C. Beddoes, W. Wallace and M.C. de Malherbe, Int. J. of Powder Metallurgy, 28, No. 3 (1992), 313. 5 J.C. Beddoes, L. Zhao and W. Wallace, in “Advancements in Synthesis and Processes”, ed. F.H. Froes et al., SAMPE (1992), M657. 6 L. Zhao, J.C. Beddoes and W. Wallace, Materials Research Soc. Symp., “High Temperature Ordered Intermetallic Alloys - V”, Boston, Nov. (1992). 7 R.J. Schaeffer and B.G. Kushner, in “Intelligent Processing of Materials”, ed. H.G.N. Wadley et al., The Minerals, Metals and Materials Society (1990), 55. 8 R.J. Schaeffer and G.M. Janowski, Acta metall. mater., 40, No 7 (1992), 1645. 9 R.J. Schaeffer, Int. J. of Powder Metallurgy, 28, No 3 (1992), 161. 10C.F. Yolton, in “P/M in Aerospace and Defence Technologies”, ed. F.H. Froes, Metal Powder Ind. Fed., Princeton, N.J., (1989), 123. 11 A.S. Helle, K.E. Easterling and M.F. Ashby, Acta metall, 33, No 12 (1985), 2163. 12R.L. Hewitt, W. Wallace and M.C. de Malherbe, Powder Metall. 16 (1973), 4798. 13M.J. Blackburn in “The Science, Technology and Applications of Titanium”, ed. R.I. Jaffe and N.E. Promisei, Pergamon Press, Oxford (1970), 633. 14 J.C. Beddoes. L. Zhao and W. Wallace, Scripta metall. mater., 28 (1993) 383.
99
Densification and microstructural control of near γ-TiAl intermetallic powders by HIPing W. Wallace*, L. Zhao* J.C. Beddoesk and D. Morphy* aStructures and Materials Laboratory, Institute for Aerospace Research, National Research Council of Canada, Ottawa, Ontario, K1A 0R6 ^Department of Mechanical and Aerospace Engineering, University, Ottawa, Ontario, K1A 5B6
Carleton
A bstract
Hot isostatic pressing has been used to consolidate pre-alloyed powders of near γ-titanium aluminide with attention being paid to the effects of HIP parameters on densification, microstructural control and properties. Three alloys have been studied, namely Ti-48A1, Ti-48A1-2W, and Ti-47.5Al-3Cr. (atom %). The W containing alloy required a temperature of 1100°C to eliminate porosity, but full density was achieved at lower temperatures for the other two alloys. While the consolidated Cr powder possessed a uniform microstructure, the other two alloys retained a dendritic structure even when HIPed at 1250°C. HIP cycles with varying temperatures and pressures were therefore designed to refine and homogenize the microstructures. The densities and microstructures resulting from the modified HIP cycles were characterized and compared to those obtained at constant temperatures and pressures. Methods of optimizing the HIP parameters to produce a homogenous microstructure suitable for further thermo-mechanical processing (TMP) or service application are discussed. 1. INTRODUCTION
The popular appeal of hot isostatic pressing (HIP) lies in its ability to heal porosity or similar defects in metal or ceramic parts or preforms. However the full potential of HIP lies in its ability to produce parts having microstructures tailored to meet the design requirements of parts in an optimum manner. This has been demonstrated at NRC and elsewhere for superalloy powder compacts and castings (1-3). Similar work, directed at the optimum HIP processing of TiAl is being carried out at NRC (4-6), and by other groups such as that of Schaeffer (7-9). In this paper the effects of HIP on the microstructure and properties of three near γ-TiAl powders are examined with particular attention to the mechanisms of densification and suitability of the compacts for their intended applications.
100
2. EXPERIMENTAL DETAILS
The powders were prepared by Crucible Compaction Metals Inc. (Oakdale, Pennsylvania) using the titanium gas atomization process (TGA), (10). The actual chemical compositions of the binary and ternary powders are Ti-48.1A1, Ti-48.4Al-1.9W and Ti-47.1Al-2.9Cr, respectively. The powders were HIP’ed in evacuated, out-gassed and sealed stainless steel or titanium containers (4). Isothermal hot compression tests were carried out on the ternary alloys in the temperature range 900-1200°C and at constant true strain rates in the range 10 4sec-i and lO -isec1. Ambient and elevated temperature tensile tests were also carried out to determine flow and fracture properties. Three types of HIP cycle were employed in the study: a) Cycle type 1: Ramping to constant temperature and pressure periods of 120 minutes, followed by in-autoclave cooling. Maximum temperature was in the range 1050°C and 1250°C, and peak pressure was 200 MPa. b) Cycle type 2: Ramping to 1380°C/200 MPa/30 minutes followed by 1100°C/170 MPa/90 minutes and in-autoclave cooling. c) Cycle type 3: Ramping to 1100°C/200 MPa/110 minutes followed by ramping to peak temperatures of 1400°C and 1415°C/200 MPa/10 minutes. Treatments (b) and (c) have only been performed on the W-containing alloy. All three treatments involved a total time of 2 hours at the hold temperatures. 3. RESULTS AND DISCUSSION 3 .1 A s-r ece iv ed p ow d ers
The binary and the Ti-48A1-2W powders had a well defined dendritic structure, whereas the Cr-containing powder was far more uniform in microstructure. The two ternary powders are compared in Figure 1, where it can be seen that light etching dendrite cores are clearly visible and quite abundant in the W-alloy. These cores are W-rich, hard and extremely stable. They resist deformation and cause plastic flow during HIP or subsequent tensile or compressive straining to be concentrated in the softer interdendritic regions. This coring is likely to inhibit plasticity and hence densification during HIP, and restrict ductility and superplasticity during subsequent tensile deformation or forging of the compacts. The high temperature treatments of HIP cycles 1 and 2 were attempts to eliminate this dendritic segregation and improve flow and fracture properties. The principal phase in the Ti-48A1 powder was γ -TiAl with a small amount (-15%) of a2. The principal phase in the W and Cr containing powders was 0C2 , with about 40% γ and about 10%βο, as illustrated for the W-containing alloy in Figure 2.
101
Figure 1. Microstructures of as-received powders, a) Ti-48A1-2W and b) Ti-47.5A1-Cr. 0.50 021 a2
As-Received Ti-48A4-2W Powder
CO
b 111-v
,
0020C 2
m 025
CL
O
,
020a2
10b
02 2 ,
L 35
40
" 7
002ώ 00γ 45 2Θ
50
55
Figure 2. X-ray diffraction plot of Ti-48A1-2W as-received powder. 3 .2 HIP c y c le s ty p e 1
The HIP temperatures were such that all alloys should be in an (X2 +J (1050 or 1100°C) or α+γ (1150° or 1250°C) condition. Full density (>99.9%) was achieved in the binary alloy and the Cr-alloy at the lowest HIP temperature employed, whereas a temperature of 1100°C was needed for the W-containing alloy. The microstructures of these two ternary materials after HIPing at 1050°C/200MPa are shown in Figure 3. Prior particle boundaries (PPB’s) can be seen in both compacts, but while the W-alloy shows an unrecrystallized structure with strong dendritic segregation, the Cr-containing alloy shows extensive recrystallization with profuse twinning. PPB’s in these alloys tend to provide paths for easy crack propagation, and hence they may impede ductility. The higher HIP temperatures, up to 1250°C can partially eliminate the PPB’s and improve homogeneity and
102
Figure 3. As-HIP 1050°C/200MPa, a) Ti-48A1-2W, b) Ti-47.5Al-3Cr.
A
Figure 4. As-HIP 1250°C/200MPa, a) Ti-48A1-2W, b) Ti-47.5Al-3Cr. ductility. However, they do not completely eliminate dendritic segregation in the W-containing alloy. The best combination of room temperature strength and ductility in the as-pressed W-alloy was 505MPa and 1.3% after 1250°C HIP (5). The microstructures of the alloys after HIPing at 1250°C/200MPa are shown in Figure 4. This higher temperature promotes recrystallization and increases grain size, as shown for the W-alloy in Table 1. In terms of phase content, both HIP temperatures of 1050°C and 1250°C cause the transformation of 0C2 to γ, as illustrated for the W-containing alloy by comparing Figure 5 with Figure 2. In this alloy after 1250°C HIP, the ßQ phase is found predominantly in the dendrite cores as small recrystallized grains, Figure 6a. In the Crcontaining alloy after 1250°C HIP, the β0 phase is fairly uniformly distributed throughout the compact, and is located at triple points and grain boundaries of the γ-phase, Figure 6b.
103
2.50
CO
ο
* 1.25
CO Q_
O
35
40
45 2Θ
50
55
Figure 5. X-ray diffraction plot of HIPped Ti-48A1-2W compact.
Table 1 Summary of the response of Ti-48A1-2W powder to Hot Isostatic Pressing No.
1. 2. 3. 4.
HIP Cycle
200MPa/1050°C/2hr 200MPa/1100°C/2hr 200MPa/1150°C/2hr 200MPa/1250°C/2hr
5. 200MPa/1380°C/0.5hr +170MPa/1100°C/1. 5hr 6. 200MPa/1100°C/ Ihr 50 min +200 MPa/ 1400°C/10 min 7. 200MPa/1100°C/lhr 50 min +200MPa/1415°C /10 min
Prior Dendritic Structure
Microstructural Average Grain Densification Size (μ) Homogeneity
retained retained retained retained
poor poor poor improved
4.1 6.0 8.0 12.5
>99.5% >99.9% >99.9% >99.9%
removed
improved
--
>99.9%
removed
removed
~200μ/Ι>25μ/γ Good (L+γ) L=ot2 +Y lam ellar Good (near fully L)
~300μ/Ι>
>99.9%
>99.9%
While mechanical properties are improved by HIPing at higher temperatures, it is unlikely that these compacts would be used for engineering components because of their low ductility. Isothermal forging is being explored as a means of refining microstructures and achieving a more favourable combination of properties. Forging preforms are therefore required that have low flow stress, and good ductility. In this case the use of high HIP temperatures is deleterious since it reduces the range of temperatures and strain rates over which superplasticity is observed, as shown in Figure 7. In this figure the contours show lines of constant strain rate sensitivity exponent m, obtained from the slope of
104
Figure 6. SEM images of ß 0 phase in a) Ti-48A1-2W and b) Ti-47.5Al-3Cr as-HIP 1250°C/200MPa. plots of log stress (σ) versus log strain rate (έ) at constant strain and temperature, according to the relationship σ = dm 3 .3 D en sifica tio n
The rates of densification of TiAl powders have been studied experimentally by Schaeffer et al. (9) using an eddy current sensor which surrounds the shrinking container. The present results are qualitatively similar to those of Schaeffer and are consistent with the models of Helle et al. (11). Initial densification occurs rapidly by plastic flow due to stress
TEST TEMPERATURE (°C)
TEST TEMPERATURE (<>C)
Figure 7. Plots of constant strain rate sensitivity exponent for Ti-48A1-2W HIPped at a)1050°C/200MPa and b) 1250°C/200MPa.
105
I
F ig u re 8 . A s-H IP ped Ti-48A1-2W , 1 1 0 0 ° C /2 0 0 M Pa. c o n c e n tr a tio n s a t p o in ts of c o n ta c t b e tw e e n p a r tic le s , a s d e s c rib e d b y H e w itt e t al. (12). E v id e n c e of p la s tic flow is fo u n d a t th e lo w e r H IP te m p e r a tu r e s in th e fo rm of d e fo rm a tio n tw in s, a n d d is lo c a tio n s e g m e n ts a n d ta n g le s a s sh o w n in Fig. 8 . T h e se fe a tu re s a re fo u n d p re d o m in a n tly in th e s o ft in te r d e n d r itic re g io n s of th e W -alloy w h e re a s th e h a r d d e n d r ite co res re m a in re la tiv e ly u n d e fo rm e d . A s th e H IP te m p e r a tu r e is in c re a s e d to 12 5 0 °C , d y n a m ic r e c r y s ta lliz a tio n o c c u rs th r o u g h o u t th e c o m p a c ts in c lu d in g th e d e n d rite co res, a n d th is e lim in a te s m o s t of th e d is lo c a tio n s . T ab le 2 E s tim a te s o f th e d e n s ity a c h ie v e d b y p la s tic y ie ld in g in T i-48A 1-2W p o w d e r a s a fu n c tio n o f HIP te m p e r a tu re In itia l D e n sity (Do)
H IP C o n d itio n s °C /M P a(P )
F low S tr e s s (ay), M P a
Yield in d u c e d d e n s ity (Dy)
0 .6 4 0 .6 4 0 .6 4
1 0 5 0 /2 0 0 1 2 5 0 /2 0 0 1 4 0 0 /2 0 0
350 150 80
0.86
0 .7 5 0 .9 8 5
C a lc u la te d from H elle e t al (11).
T h e e q u a tio n s of H elle e t al. (11) c a n b e u s e d to c a lc u la te th e in itia l d e n s if ic a tio n t h a t o c c u r s b y p la s tic y ie ld in g , a n d t h i s p ro v id e s a n in d ic a tio n o f th e r e s id u a l p o ro s ity t h a t m u s t b e e lim in a te d b y c re e p o r o th e r m e c h a n is m s . T h e r e s u lts of th e s e c a lc u la tio n s a re s u m m a riz e d in
106
Table 2, where it can be seen that even a temperature as high as 1400°C would not achieve a density of 99% by plastic flow alone. In these calculations the initial density of the compacts was taken to be 0.64, the initial strain rate was estimated from the results of Schaeffer to be about 2.7 x 10-2 s-i and values of flow stress were obtained from the true stress versus true strain compression tests at the appropriate temperature and strain rate The residual porosity to be closed by creep or diffusion therefore varies from about 0.25 at 1050°C to about 0.015 at 1400°C, and evidence of creep is seen in Figure 8 in the form of dislocation sub-boundaries (marked “S”). A further mechanism of densification to consider is that of transformation induced plasticity, as proposed by Schaeffer and Janowski (8). They have suggested that in powders consisting predominantly of 2 phase, the transformation to γ will result in the formation of metastable 0C stacking faults and the passage of dislocations through the 0C2 phase, which contribute to densification by local deformation in the region of powder particle contact. The crystallographic relationships maintained by this transformation are those determined by Blackburn (13), such that the two phases share close packed planes and directions, (0002)α2//(111)γ and <1120>α2//<110>γ. As illustrated by comparing Figures 2 and 5, a substantial transformation of 0C2 to γ occurs in the W-containing powder, and therefore transformation induced plasticity as described by Schaeffer and Janowski is a possible contributor to densification in this alloy. 3 .4 HIP c y c le ty p e 2
The purpose of this treatment was to eliminate dendritic segregation while allowing only limited grain growth. The resulting microstructure contained large areas of equiaxed and fine-grained phases owing to
Figure 9. Ti-48A1-2W a) cycle 2, b) cycle 3 to 1400°C, c) cycle 3 to 1415°C.
107
complete recrystallization. Moreover, the dendritic structure no longer existed indicating that segregation was unstable under this HIP condition. While the overall microstructural homogenity was improved, regions containing coarse grains of γ-phase existed, Fig. 9a (arrows), with a typical size close to that of prior particles. This was attributed to the strong dendritic segregation in some powder particles where the interdendritic region dominated by γ single phase coarsened rapidly during HIPing due to the lack of second phase (4). Mechanical properties have not been determined for this structure. 3 .5 HIP c y c le ty p e 3
In these cycles the two stage process of cycle 2 was reversed to achieve initial densification and homogenization in the α2+γ phase field at 1100°C before completing the transformation to a at 1400°C. A thermocouple was mounted on the HIP can, and the can was packed in sand to ensure a slow heat up to the peak temperature. The microstructure, Figure 9b, consisted of a mixture of fine, equiaxed grained γ, and coarser grained lamellar α2+γ structure of the type desired. For this composition (Ti-48A12W), 1400°C is just above the a transus (14), but it appears that a fully transformed a structure was not achieved. In the latest run, the peak temperature was raised to 1415°C but the can was not packed in sand. Figure 9c shows that an almost fully transformed lamellar structure of α2+γ was obtained (>90%) but a small amount of equiaxed γ was retained. It is possible that the material did not experience the desired temperature of 1415°C for sufficiently long time to complete the transformation. 4.
FUTURE WORK
This work is continuing to produce a fully transformed α2+γ structure at which time the mechanical properties of the as-HIP material will be determined. Work is also continuing on the isothermal forging response of the Ti-48Al-3Cr alloy, which will determine the envelope for superplastic forging and elucidate the role of the β0 phase, Fig. 6b. 5.
CONCLUSIONS
• The presence of W in TiAl powders promotes dendritic segregation which is stable and impedes densification by plastic yielding. Chromium additions provide greater homogeneity and enhanced densification. • Densification occurs by plastic yielding, creep and transformation induced plasticity (α2->γ) in the W alloy, and probably the Cr alloy.
108
• High HIP temperatures in the α+γ phase field achieve full density and improved strength and ductility, but they reduce the envelope for superplastic flow during forging. • Even higher HIP temperatures in the single phase a field are feasible and should allow the formation of fully transformed α2+γ lamellar structures. 6 . ACKNOWLEDGEMENTS
This work was supported by the National Research Council of Canada under IAR-SML project JHR-01. Financial assistance was also provided by the Department of National Defence via Financial Arrangement 220791NRC51. The support of both organizations is gratefully acknowledged. 7 . REFERENCES
1 W. Wallace, J-P. Immarigeon and J.M. Trenouth, AGARD C.P. 200, Neuilly-sur-Seine (1976). 2 W. Wallace, H.B. Dunthome and R. Sprague, Can. Met. Quart., 13, No. 3 (1974), 517. 3 J.C. Beddoes and W. Wallace, Metallography, 13 (1980) 185. 4 J.C. Beddoes, W. Wallace and M.C. de Malherbe, Int. J. of Powder Metallurgy, 28, No. 3 (1992), 313. 5 J.C. Beddoes, L. Zhao and W. Wallace, in “Advancements in Synthesis and Processes”, ed. F.H. Froes et al., SAMPE (1992), M657. 6 L. Zhao, J.C. Beddoes and W. Wallace, Materials Research Soc. Symp., “High Temperature Ordered Intermetallic Alloys - V”, Boston, Nov. (1992). 7 R.J. Schaeffer and B.G. Kushner, in “Intelligent Processing of Materials”, ed. H.G.N. Wadley et al., The Minerals, Metals and Materials Society (1990), 55. 8 R.J. Schaeffer and G.M. Janowski, Acta metall. mater., 40, No 7 (1992), 1645. 9 R.J. Schaeffer, Int. J. of Powder Metallurgy, 28, No 3 (1992), 161. 10C.F. Yolton, in “P/M in Aerospace and Defence Technologies”, ed. F.H. Froes, Metal Powder Ind. Fed., Princeton, N.J., (1989), 123. 11 A.S. Helle, K.E. Easterling and M.F. Ashby, Acta metall, 33, No 12 (1985), 2163. 12R.L. Hewitt, W. Wallace and M.C. de Malherbe, Powder Metall. 16 (1973), 4798. 13M.J. Blackburn in “The Science, Technology and Applications of Titanium”, ed. R.I. Jaffe and N.E. Promisei, Pergamon Press, Oxford (1970), 633. 14 J.C. Beddoes. L. Zhao and W. Wallace, Scripta metall. mater., 28 (1993) 383.