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Effects of strain rate and temperature on deformation behaviour of IN 718 during high temperature deformation
Materials' Science and Engineering, A 177 (1994) 1-9
1
Effects of strain rate and temperature on deformation behaviour of IN 718 during high temperature deformation L. X. Z h o u a n d T. N. B a k e r
Department of Metallurgy and Engineering Materials, University of Strathclyde, Glasgow G 1 1XN (UK) (Received March 12, 1993)
Abstract The hot deformation characteristics of a wrought IN 718 alloy were investigated by compression testing at constant strain rates in the range of 0.1 to 5 × 10 3 s- ~, and testing temperatures in the range of 950 to 1100 °C using a 200 ton capacity microprocessor controlled Fielding hydraulic press. Examination of the microstructures was carried out by optical microscopy and TEM. The flow stress of the compression tests showed a single peak in the flow stress-strain curves, and indicated that a dynamic recrystallization transition took place during the hot compression. The relationship between the peak stresses (up) and the Zener-Hollomon parameter (z) can be expressed by G = 0'5Z° ~7. "Necklace" microstructures were observed at testing temperatures below 1050 °C, for strain of 0.7. The fraction of recrystallized grains increased with the increasing temperature and strain, and decreasing strain rate. Fully recrystallized microstructures were observed at temperatures 1050 °C or greater, with a strain of 0.7.
1. Introduction The behaviour of metals and alloys during high temperature deformation is complex and changes with the changing processing parameters, i.e. strain, strain rate and temperature. High temperature deformation is associated with dynamic restoration processes which affect the microstructure and properties. It is important to understand the mechanisms of high temperature deformation and the relationship between the processing variables, microstructure, and properties. Since the late 1960s, extensive research on these relationships has been carried out on steel and nickel alloys [1-8], and copper [9-10], theoretically and experimentally. However, little has been published either in the open literature or in reports on IN718 alloy. One of the reasons for this is the complex chemical composition of IN 718 alloy, and especially the intermetallic phase transformation which occurs during the high temperature deformation. IN 718 alloy is nickel rich, and strengthened predominately by ordered gamma double prime (y") precipitates. Because of its excellent mechanical properties up to a temperature as high as 650 °C, the alloy has emerged as one of the most widely used superalloys of recent times. Over the temperature range approximately 650 to 1000 °C there is a complex precipitation sequence [Ni3(A1,Ti, Nb ) 7'fcc_bct~Ni3Nb 7b¢~Ni3 Nb 6orth .... bic] with increasing temperature. Thus 6 solvus temperature is 1000 °C. 0921-5093/94/$7.00 SSD1 0921-5(/93(93)09402-5
The present study was aimed at investigating the effect on thermo-mechanical behaviour and microstructure, of strain rates 0.1, 5 × 10 -2, and 5 × 10 -3 s- ~over a temperature range from 950 to 1100 °C, and understanding the relationship between dynamic restoration processes and the flow stress-strain curves. Changes of microstructure associated with different degrees of reduction during the compression processing, were also investigated.
2. Experimental materials and procedure The chemical composition of IN 718 alloy used in this investigation is given in Table 1. Cylindrical samples, 34 mm in diameter and 45 mm in height, were prepared from an extruded 200 mm diameter remelted billet. The microstructures of the starting material showed that it was uniform and 6 phase free, and the initial grain size was 114/~m (ASTM 3). The high temperature compression tests were carried out on a 200 ton capacity microprocessor controlled Fielding hydraulic press, a testing system which is designed for constant strain rate compression under high temperature conditions. The arrangement of the specimen, insulated and dies is shown in Fig. 1. In order to achieve accurate data of true stresses, pressing forces were measured by a load cell instead of the pressure meter of press. The load cell was located in the press© 1994 - Elsevier Sequoia. All rights reserved
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L.X. Zhou, T. N. Baker
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Behaviour oflN 718 during high temperature deformation
ing system, as shown in Fig. 1, so that frictional effects on the measurement were avoided. The specimens were induction heated up to testing temperatures and soaked at temperature for 20 min before compressing. Tests were conducted over a temperature range of 950 to 1100 °C at true strain rates of 0.1, 5 x 10- 2 and 5 x 10- 3 s- 1 followed immediately by water quenching, with a delay of less than 2 s, so as to preserve the hot deformed structures. All the specimens were deformed by 50% reduction, except those used for the examination of microstructural changes during the compression process, which were quenched after different strains. The data of time-load-displacement were recorded by a microprocessor system at intervals of 0.1 s. TABLE 1. Chemical composition of the material tested Element
Content (wt.%)
A1 B C Co Cr Fe Mo Nb Ni Ti
0.06 0.004 0.037 0.16 18.35 17.64 2.98 5.11 53.88 0.96
3. R e s u l t s
3.1 Stress-strain behaviour Typical flow curves for IN 718 alloys are presented in Fig. 2. The flow stress curves are for a testing temperature range of 950 to l l00°C, deformed at constant strain rates of 0.1 s -1 and 5 × 10 -2 s -1. The characteristics of the curves in Fig. 2 are (i) a single peak curve, (ii) the rapid initial rise to a peak stress, followed by (iii) a gradually decreasing flow stress to around 0.7 strain. These features represent the relationship between flow stress and strain associated with dynamic recovery combined with dynamic recrystallization [2, 4, 11]. It is seen that the peak stresses increase with decreasing temperature and the higher the temperature, the less is the difference in the flow stress developed by same testing temperature interval. The strains of the peak stresses decreased with increasing temperature. The strain rate dependence of flow stress is illustrated in Fig. 3. The flow stress is mildly sensitive to the strain rates employed. The peak stresses increase with the increasing strain rates at a fixed testing temperature. The shape of the curves shows a slight change, and it becomes flatter as the strain rate is reduced. The strain rate sensitivity of flow stress of IN 718 alloy at 1000 and 1050 °C was determined to be 0.156 and 0.166 respectively. The strain rate sensitivity is also temperature dependent, and increased with temperature.
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/ Behaviour of IN 718 during high temperature de#ormation
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Fig. 2. Effect of temperature on flow stress-strain curves in compression testing of IN 718 alloy at strain rate of (a) 0.1 s-1 and ( b ) 5 x l 0 Zs ].
(b)
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Fig. 3. Strain rate dependence of the flow stress tested at (a) 1000 °C and (b) 1050 °C.
4
L.X. Zhou, T. N. Baker
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Behaviour of IN 718 during high temperature deformation
3.2. Metallography
No 6 phase was found in the specimen deformed between 1000 and l l 0 0 ° C at strain rates of 0.1, 5 x 10 2 and 5 x 10 -3 s -1. A negligible amount of 6 phase was observed in the microstructure developed at the temperature of 950 °C. An example of 6 phase is shown in Fig. 4. Dynamic recovery and dynamic recrystallization were observed in the microstructure throughout the range of the experiments. Typical microstructures are presented in Fig. 5, in which can be seen the range of different dislocation densities commonly observed under these conditions, and associated with critically worked, partly worked, as well as nearly dislocation free regions. The dislocation densities in dynamically recrystallized grains are different, as shown in Fig. 5(b). The nucleation of a dynamic recrystallized grain at the peak strain ep, is shown in Fig. 6(a). It also can be seen in Fig. 5(a) that the dislocation density in the deformed
The microstructure was examined using optical and transmission electron microscopy. The specimens for optical microscope were etched by using a solution of 50% hydrogen peroxide, 45% hydrochloric acid, and 5% hydrofluoric acid. The foils for TEM were jet polished at - 35 to - 40 °C in a 10% perchloric acid, 30% n-butanol and 60% ethanol solution [12]. Non-uniform deformation was observed in the regions which were near top and bottom of an asforged specimen. This is influenced by the friction. During deformation, if the strain rate is too high, because of either the ram speed of the forging equipment, or localized preferential flow within the forging, then a local temperature increase can occur by adiabatic heating. In the present study, the strain rate range is relatively low, so that adiabatic heating can be neglected.
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Fig. 4. 6 phase produced at temperature of 950 °C, and strain rate of 0.1 s- 1: (a) bright field and dark field images; (b) the EDX spectrum from the 6 particle region; (c) selected area electron diffraction pattern; and (d) indexed pattern of matrix ), and particle 6 which give an orientation relationship [110]~ ]1[10016. o matrix 7,, • particle 6.
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Behaviour of lN 718 during high temperature deformation 350
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Fig. 5. Typical substructure of dynamically recrystallized grains deformed at (a) 1000 °C and (b) 1050 °C. region is considerable. The effect of strain and deformation temperature on the progress of dynamic recrystallization are illustrated in Figs. 5 and 6. For a given strain of around 0.7, the fraction of recrystallized grains increased with increasing temperature, as shown in Fig. 7. The necklace structure in Fig. 6(b) shows that the preferred nucleation sites of recrystallized grains are at strained grain boundaries, which is a significant feature of the recrystallized structures of IN 718 alloy. The fully recrystallized microstructures were obtained at strain of around 0.7 at temperature of 1050 °C and above, as seen in Fig. 6(c). The fraction of recrystallized grains increased with strain rate in the range of 0.1 to 5 × 10 -3 s -l, but the influence of strain rate on this fraction is much less than the influence of temperature.
4. Discussion 4.1. Flow stress curve behaviour a n d d y n a m i c recrystallization
The flow stress curves of IN 718 alloy developed under high temperature compression conditions, shown in Fig. 2, belong to those of the "normal" be-
Fig. 6. True stress-strain curve and microstructure evolution due to increasing strain, tested at 1050 °C and 0.1 s-~ strain rate: (a) initiation of dynamic recrystallization at the grain boundaries at e = ep; (b) typical necklace structure obtained at ~'= 0.3; (c) fully dynamically recrystallized microstructure achieved at e = 0.75.
haviour of the alloy group with a low or intermediate value of stacking fault energy, and are similar to the curves for nickel [1, 13, 14], copper and austenitic steels [1, 4, 15, 16]. The initial rapid rise in stress is associated with an increase in dislocation density and the formation of poorly developed subgrain boundaries, as a result of work hardening and dynamic recovery. In alloys with a low or intermediate stacking fault energy, e.g. nickel or stainless steel, the dynamic recovery process is very slow, and the dislocation density can attain a sufficiently high value for dynamic recrystallization to be initiated, once a critical strain is exceeded. The role played by the dislocation density in the initiation of dynamic recrystallization has been examined previously [11, 17, 18]. It is apparent that a sufficient energy density must be built up to provide a driving force for the growth of nuclei. It can be seen that ep increased as the strain rate increased and the temperature decreased, as represented in Figs. 2 and 3. The recrystallized grains were observed at ep, as shown in Fig. 6(a) which indicates
6
L.X. Zhou, T. N. Baker
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Behaviour oflN 718 during high temperature deformation
b the Burgers vector; r the dislocation line energy; dsb the subgrain diameter; m the mobility; and g the strain rate. Hence a critical dislocation density, which depends on the deformation conditions and the grain-boundary characteristics (grain-boundary energy and mobility), must be exceeded before a stable nucleus for dynamic recrystallization can be formed. As a rule, the grain-boundary mobility conforms to an Arrhenius law [22], i.e. m = m 0 exp R T -
Fig. 7. Dynamic recrystallizationinfluenced by the temperature tested at g = 0.1 s-1 and (a) 950 °C, (b) 1050 °C.
that dynamic recrystallization started before peak strain was reached. Under the conditions of strain, nucleation originates at pre-existing grain boundaries, where the dislocation density is sufficient to initiate dynamic recrystallization. The nucleation of recrystallized grains has been found at around 0.5 to 0.7 ep in earlier work [7, 19, 20, 21], and it can be assumed that the critical strain to initiate dynamic recrystallization is equal to •p [19]. It is further indicated that a critical dislocation density, which depends on the conditions of deformation, must be exceeded before recrystallization can proceed. The effects of temperature on the critical strain are similar to those of e_. This can be explained by the analysis of Roberts Iv4], where the dynamic recrystallization occurs at a critical value of dislocation density which is expressed as
p0-const./be,7
) =po
(1>
where P0 is the dislocation density in unrecrystallized material; P0c the critical dislocation density for dynamic recrystallization; s the grain-boundary energy;
(2)
where m0 is a constant. Qm is an activation energy. T and R are, respectively, temperature and gas constant. An increasing temperature leads to an increase in grain-boundary mobility and a decrease in the critical dislocation density for dynamic recrystallization, which in turn results in a decrease in the peak strain ep. Strain rate has the opposite effect. The driving force for the growth of dynamically recrystallized grains is gradient of dislocation density. Figure 5 shows typical features of dynamic recrystallization. A new grain at (R) on Fig. 5(a) grows by consuming adjacent cell structures which contain a much higher dislocation density. In the new grain, dislocations are present, indicating that this microstructure has been developed by dynamic recrystallization. This is because dynamically recrystallized grain growth is accompanied by deformation processing which strains the new grain and increases the dislocations in it. As a result of this current deformation, the dislocation densities in dynamically recrystallized grains are various and change with their extent of strain, as shown in Fig. 5(b). The fraction of dynamically recrystallized grains increases with the temperature, and the fully recrystallized microstructure was obtained at a temperature greater than 1050°C and strain of around 0.7, as shown in Fig. 7. Because the dislocations have increased mobility through cross-slip and climb, this permits attractive interactions to cause annihilation, and a rearrangement into a lower energy, more regular, substructure. Microstructural change with increasing strain is shown in Fig. 6. The progress of dynamic recrystallization is developed by increasing the strain after initiation at ep. The dynamically recrystallized grains are equiaxed with a low dislocation density induced by concurrent deformation. Increasing the strain generates more dislocations in the alloy, which creates a greater driving force for nucleation and grain growth. The absence of 6 phase in the range 1000 to 1100 °C for all strain rate conditions used in this work,
L. X. Zhou, T. N. Baker
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Behaviour of IN 718 during high temperature de&rmation
illustrates that the solvus temperature of 6 phase is not affected by the strain rate of the starting material with a large grain size and which is 6 phase free. However, for a testing temperature below 6 solvus temperature (1000°C), 6 phase precipitates during deformation, and usually at grain boundaries, as shown in Fig. 4. 4.2. 7he Zener-Hollomon parameter Z The peak stress increases with increasing temperature and decreasing strain rate in Figs. 2 and 3. The data for peak stress, strain to peak stress and Z, obtained for IN 718 alloy, are given in Table 2. The correlation of the peak flow stress ap with temperature T and strain rate ~ during hot working can be expressed by the following mechanical equation of state:
%=AZ"
(3)
or
o,,=Ae"
{,,01
(4)
where A and n are constants. The flow stress increase with strain rate and the effect at constant strain can be expressed by eqn. (5), and the temperature dependence of flow stress can be expressed by eqn. (6).
Equation (4) consists of both following equations:
o=c "
(5)
A(a)=~ee/kr
(6)
where C is a strength constant which depends upon strain, temperature, and alloy composition, and Q is the activation energy for hot working. R is the gas constant, and n' is the strain rate sensitivity of the flow stress. As shown in Fig. 8, eqn. (3) agrees very well with the experimental data. n was determined to be 0.17 _+0.015, which is similar to the data for strain rate sensitivity calculated from the flow strain curves at temperatures 1000 and 1050 °C. The value of Q for IN 718 alloy can be determined from Fig. 8, as Q = 400 +_25 kJ mol-~, in agreement with the results of Guimaraes and Jonas [13] and Camus et al. [14]. The activation energies obtained in hot torsion, hot compression and creep for IN 718 alloy, nickel and nickel base alloys are listed in Table 3. There is a good agreement in the data for Q for IN 718 alloy, and this
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T A B L E 2. Observed values of ap and 0o.7 u n d e r various Z condition
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Temp.
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(°C)
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(MPa)
1000 1050 950 1000 1050 950 1000 1050
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152 120 305 200 150 317 225 175 148
112 82 225 148 110 -168 125
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1110
100
101 1013
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Fig. 8. Correlation of peak stress with Z e n e r - H o l l o m o n parameter for IN 718 alloy in present work ( ) and previous work conducted by hot compression from ref. 13 (. . . . ) and by hot torsion from ref" 14 ( ..... ).
T A B L E 3. Comparative values of n and activation energy Q Material
Process
Q (kJ tool- 1)
Reference
IN 718 IN 718 IN 718 Ni Ee- 1 8 % C r - 8 % N i Fe- 1 8 % C r - 11%Ni Fe- 1 8 % C r - 1 l%Ni
Hot compression Hot torsion Hot compression Hot working Hot torsion Hot torsion Creep
400 + 25 400 _+ 15 400 + 50 297 414 410 314
Present work [ 14] [ 13] [25] [ 1] [ 15] [ 16]
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L.X. Zhou, T. N. Baker
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BehaviouroflN 718duringhigh temperaturedeformation
implies that the activation energy, Q, is not affected by the experimental mode. For nickel base alloys, the values derived from high strain rate tests are much higher than those for creep, because dynamic restoration was by recrystallization in the former and by polygonization in the latter. The value of Q for the pure metal is, as expected, much lower than that for alloys.
4.3. Effect o f strain rate on microstructure
The effect of strain rate in the range from 0.1 to 5 x 10-3 s-1 on the microstructure is not very strong, which is in good agreement with the conclusion reported by Barker et al. [23]. The results are different from those published by Camus et al. in 1986 [14]. There are two reasons for the difference. First, the experiments of Camus et al. were conducted by hot torsion, which is a different mode of testing from the high temperature compression test employed in the present research. Also the flow stress-strain data was translated from the torsion mode to a true stress-strain. In the present work, using the system designed for high temperature compression, the true stress and strain were obtained directly from the recorded data. Comparing the test results from torsion with those obtained from compression, it is found that the torsion data is somewhat higher than the equivalent quantity measured in compression, and similar to that reported previously [24]. Additionally, the range of strain rate in the two studies was different. The work of Camus et al. [14] was carried out using a higher strain rate range from 4 to 4 × 10 -2 s -1, compared with the strain rate range used in the present work which was from 0.1 to 5 x 10- 3 s- i. Also in the work of Camus et al. [ 14], a large amount of metadynamic recrystallization was observed in the specimens produced at the higher strain. Metadynamic recrystallization is a zero incubation recrystallization, based on the existing nuclei which proceed recrystallization and cause an increase in grain growth compared with dynamic recrystallization, because there is no concurrent deformation to restrict grain growth. The strain rate has a very strong effect on the metadynamic recrystallization progress, especially on the delay time (the time from the deformation process ceasing to the metadynamic recrystallization starting); the higher the strain rate, the shorter the delay time. Metadynamic recrystallization could start after less than 1 s at a strain rate as high as 10 s- 1. Because there is some metadynamic recrystallization found in this work, the grain size developed under the higher strain rates is not due exclusively to the dynamic recrystallization is greater than that formed by dynamic recrystallization alone.
o
Conclusions
(1)
In IN 718 alloy, it has been shown experimentally from compression testing in a 200 ton microprocessor controlled press, that flow stresses are temperature and strain rate dependent, and increase with decreasing temperature and increasing strain rate. (2) The relation between peak stress Op and the corresponding regime for dynamic recrystallization can be expressed by the equation Op-~- 0 . 5 Z 0.17
where Z is the Zener-Holloman parameter. (3) The strain rate sensitivities of the flow stress at 1000 and 1050°C are 0.156 and 0.166 respectively. (4) Dynamic recrystallization occurs during the deformation over a temperature range 950 to 1100 °C, and increases with the temperature. A fully dynamically recrystallized microstructure can be achieved at temperatures above 1050 °C. (5) The solvus temperature of 6 phase is not influenced by the strain rates of deformation used in this work. (6) The strain ep to peak stress Op is also affected by the temperature and strain. (7) Initiation of dynamic recrystallization starts at ep and the favourable nucleation sites are the original grain boundaries. (8) Before steady state conditions are reached, increasing strain leads to dynamic recrystallization progress. (9) The rate of metadynamic recrystallization is related to the processing parameters, and increases with temperature and strain rate.
Acknowledgments The authors wish to thank Cameron Forged Products of Livingstone for financial support. Acknowledgments are also made to the Committee of Vice-Chancellors and Principals of the University of the United Kingdom for an ORS award to one of the authors (L.X.Z.).
References 1 J.J. Jonas, C. M. Sellars and W. J. McG. Tegart, Metall. Rev., 14(1969) 1. 2 H.J. McQueen, Microstruct. Sci., 7(1979) 71. 3 C.M. Sellars, Met. Forum, 4(1981) 75.
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Behaviour o~ IN 718 during high temperature deformation
4 W. Roberts, in G. Krauss (ed.), Deformation, Processing and Structure, AM. Soc. Metals, Metal Park, OH, 1984, pp. 109-184. 5 T. Sakai and J. J. Jonas, Acre Metall., 32 (1984), 189. 6 M.J. Luton and C. M. Sellars. Acta Metall., 17(196911033. 7 H.J. McQueen, Mater. Sci. Eng. A, 101 (19881 149. 8 T. Sakai, M. Ohashi and K. Chiba, Acta Metall., 36 (19881 1781. 9 L. Blaz, T. Sakai and J. J. Jonas, Met. Sci., 17( 19831609. 10 H.J. McQueen and S. Bergerson, Met. Sci., 6 (1972) 25. 11 J.J. Jonas, l'roc. 4th Int. Coqf Strength of Metals and Alloys, Nan~y, France, 1976, Vol. 2, pp. 976-1001. 12 J. E Collier, A. O. Selius and J. K. Tien, Superalloy 1984, AIME, p. 43. 13 A. A. Guimaraes and J. J. Jonas, Memll. Trans. A, 12 (1981 ) 1655. 14 G. Camus, B. Pieraggi and F. Chevet, Proc. Symp. co-sponsored by the mechanical metallurgy and shaping and forming committees of TMS-AIME, October 1986, pp. 305-325. 15 D.R. Barraclough and C. M. Sellars, Inst. Phys. Conf. Ser., 21, 1974.
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16 F. Garofalo, C. Richmond, W. F. Domis and F. Von Gemmingen, Joint Int. Conf. ('reep, Institute of Mechanical Engineers, London, 1963, pp. 31-39. 17 R. Sandstrom and R. Langeborg, Acta Metall., 32 (19751 387. 18 W. Roberts and B. Ahlblom, Proc. 4th Int. ('onf. Strength of Metals and Alloys, Nancy, France, 1976, Vol. 1, pp. 400-404. 19 T. Sakai and M. Ohashi, Mater. Sci. Technol., 6 (1990) 1251. 20 J. E Sah, G. J. Richardson and C. M. Sellars, Met. Sci., (19741325. 21 D. W. Livesey and C. M. Sellars, Mater. Sci. Technol., 1 (19851 136. 22 W. Roberts and B. Ahlblom, Acta Metall., 26 ( 197818(11. 23 J.F. Barker, D. D. Krueger and D. R. Chang, Proc. Advanced ttigh Temperature Alloys, ASM, Boston, USA, Oct. 1984, pp. 125-137. 24 J.J. Jonas and T. Sakai, Deformation, in G. Krauss (ed.), Processing and Structure, AM. Soc. Metals, Metal Park, OH, 1984, pp. 185-230. 25 C. M. Sellars and W. J. McG. Tegart, Mere. Sci. Rev, Metall., (~3(1966) 731.
1
Effects of strain rate and temperature on deformation behaviour of IN 718 during high temperature deformation L. X. Z h o u a n d T. N. B a k e r
Department of Metallurgy and Engineering Materials, University of Strathclyde, Glasgow G 1 1XN (UK) (Received March 12, 1993)
Abstract The hot deformation characteristics of a wrought IN 718 alloy were investigated by compression testing at constant strain rates in the range of 0.1 to 5 × 10 3 s- ~, and testing temperatures in the range of 950 to 1100 °C using a 200 ton capacity microprocessor controlled Fielding hydraulic press. Examination of the microstructures was carried out by optical microscopy and TEM. The flow stress of the compression tests showed a single peak in the flow stress-strain curves, and indicated that a dynamic recrystallization transition took place during the hot compression. The relationship between the peak stresses (up) and the Zener-Hollomon parameter (z) can be expressed by G = 0'5Z° ~7. "Necklace" microstructures were observed at testing temperatures below 1050 °C, for strain of 0.7. The fraction of recrystallized grains increased with the increasing temperature and strain, and decreasing strain rate. Fully recrystallized microstructures were observed at temperatures 1050 °C or greater, with a strain of 0.7.
1. Introduction The behaviour of metals and alloys during high temperature deformation is complex and changes with the changing processing parameters, i.e. strain, strain rate and temperature. High temperature deformation is associated with dynamic restoration processes which affect the microstructure and properties. It is important to understand the mechanisms of high temperature deformation and the relationship between the processing variables, microstructure, and properties. Since the late 1960s, extensive research on these relationships has been carried out on steel and nickel alloys [1-8], and copper [9-10], theoretically and experimentally. However, little has been published either in the open literature or in reports on IN718 alloy. One of the reasons for this is the complex chemical composition of IN 718 alloy, and especially the intermetallic phase transformation which occurs during the high temperature deformation. IN 718 alloy is nickel rich, and strengthened predominately by ordered gamma double prime (y") precipitates. Because of its excellent mechanical properties up to a temperature as high as 650 °C, the alloy has emerged as one of the most widely used superalloys of recent times. Over the temperature range approximately 650 to 1000 °C there is a complex precipitation sequence [Ni3(A1,Ti, Nb ) 7'fcc_bct~Ni3Nb 7b¢~Ni3 Nb 6orth .... bic] with increasing temperature. Thus 6 solvus temperature is 1000 °C. 0921-5093/94/$7.00 SSD1 0921-5(/93(93)09402-5
The present study was aimed at investigating the effect on thermo-mechanical behaviour and microstructure, of strain rates 0.1, 5 × 10 -2, and 5 × 10 -3 s- ~over a temperature range from 950 to 1100 °C, and understanding the relationship between dynamic restoration processes and the flow stress-strain curves. Changes of microstructure associated with different degrees of reduction during the compression processing, were also investigated.
2. Experimental materials and procedure The chemical composition of IN 718 alloy used in this investigation is given in Table 1. Cylindrical samples, 34 mm in diameter and 45 mm in height, were prepared from an extruded 200 mm diameter remelted billet. The microstructures of the starting material showed that it was uniform and 6 phase free, and the initial grain size was 114/~m (ASTM 3). The high temperature compression tests were carried out on a 200 ton capacity microprocessor controlled Fielding hydraulic press, a testing system which is designed for constant strain rate compression under high temperature conditions. The arrangement of the specimen, insulated and dies is shown in Fig. 1. In order to achieve accurate data of true stresses, pressing forces were measured by a load cell instead of the pressure meter of press. The load cell was located in the press© 1994 - Elsevier Sequoia. All rights reserved
2
L.X. Zhou, T. N. Baker
/
Behaviour oflN 718 during high temperature deformation
ing system, as shown in Fig. 1, so that frictional effects on the measurement were avoided. The specimens were induction heated up to testing temperatures and soaked at temperature for 20 min before compressing. Tests were conducted over a temperature range of 950 to 1100 °C at true strain rates of 0.1, 5 x 10- 2 and 5 x 10- 3 s- 1 followed immediately by water quenching, with a delay of less than 2 s, so as to preserve the hot deformed structures. All the specimens were deformed by 50% reduction, except those used for the examination of microstructural changes during the compression process, which were quenched after different strains. The data of time-load-displacement were recorded by a microprocessor system at intervals of 0.1 s. TABLE 1. Chemical composition of the material tested Element
Content (wt.%)
A1 B C Co Cr Fe Mo Nb Ni Ti
0.06 0.004 0.037 0.16 18.35 17.64 2.98 5.11 53.88 0.96
3. R e s u l t s
3.1 Stress-strain behaviour Typical flow curves for IN 718 alloys are presented in Fig. 2. The flow stress curves are for a testing temperature range of 950 to l l00°C, deformed at constant strain rates of 0.1 s -1 and 5 × 10 -2 s -1. The characteristics of the curves in Fig. 2 are (i) a single peak curve, (ii) the rapid initial rise to a peak stress, followed by (iii) a gradually decreasing flow stress to around 0.7 strain. These features represent the relationship between flow stress and strain associated with dynamic recovery combined with dynamic recrystallization [2, 4, 11]. It is seen that the peak stresses increase with decreasing temperature and the higher the temperature, the less is the difference in the flow stress developed by same testing temperature interval. The strains of the peak stresses decreased with increasing temperature. The strain rate dependence of flow stress is illustrated in Fig. 3. The flow stress is mildly sensitive to the strain rates employed. The peak stresses increase with the increasing strain rates at a fixed testing temperature. The shape of the curves shows a slight change, and it becomes flatter as the strain rate is reduced. The strain rate sensitivity of flow stress of IN 718 alloy at 1000 and 1050 °C was determined to be 0.156 and 0.166 respectively. The strain rate sensitivity is also temperature dependent, and increased with temperature.
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/ Behaviour of IN 718 during high temperature de#ormation
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Fig. 2. Effect of temperature on flow stress-strain curves in compression testing of IN 718 alloy at strain rate of (a) 0.1 s-1 and ( b ) 5 x l 0 Zs ].
(b)
=
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Fig. 3. Strain rate dependence of the flow stress tested at (a) 1000 °C and (b) 1050 °C.
4
L.X. Zhou, T. N. Baker
/
Behaviour of IN 718 during high temperature deformation
3.2. Metallography
No 6 phase was found in the specimen deformed between 1000 and l l 0 0 ° C at strain rates of 0.1, 5 x 10 2 and 5 x 10 -3 s -1. A negligible amount of 6 phase was observed in the microstructure developed at the temperature of 950 °C. An example of 6 phase is shown in Fig. 4. Dynamic recovery and dynamic recrystallization were observed in the microstructure throughout the range of the experiments. Typical microstructures are presented in Fig. 5, in which can be seen the range of different dislocation densities commonly observed under these conditions, and associated with critically worked, partly worked, as well as nearly dislocation free regions. The dislocation densities in dynamically recrystallized grains are different, as shown in Fig. 5(b). The nucleation of a dynamic recrystallized grain at the peak strain ep, is shown in Fig. 6(a). It also can be seen in Fig. 5(a) that the dislocation density in the deformed
The microstructure was examined using optical and transmission electron microscopy. The specimens for optical microscope were etched by using a solution of 50% hydrogen peroxide, 45% hydrochloric acid, and 5% hydrofluoric acid. The foils for TEM were jet polished at - 35 to - 40 °C in a 10% perchloric acid, 30% n-butanol and 60% ethanol solution [12]. Non-uniform deformation was observed in the regions which were near top and bottom of an asforged specimen. This is influenced by the friction. During deformation, if the strain rate is too high, because of either the ram speed of the forging equipment, or localized preferential flow within the forging, then a local temperature increase can occur by adiabatic heating. In the present study, the strain rate range is relatively low, so that adiabatic heating can be neglected.
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Fig. 4. 6 phase produced at temperature of 950 °C, and strain rate of 0.1 s- 1: (a) bright field and dark field images; (b) the EDX spectrum from the 6 particle region; (c) selected area electron diffraction pattern; and (d) indexed pattern of matrix ), and particle 6 which give an orientation relationship [110]~ ]1[10016. o matrix 7,, • particle 6.
L. X. Zhou, T. N. Baker
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Behaviour of lN 718 during high temperature deformation 350
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Fig. 5. Typical substructure of dynamically recrystallized grains deformed at (a) 1000 °C and (b) 1050 °C. region is considerable. The effect of strain and deformation temperature on the progress of dynamic recrystallization are illustrated in Figs. 5 and 6. For a given strain of around 0.7, the fraction of recrystallized grains increased with increasing temperature, as shown in Fig. 7. The necklace structure in Fig. 6(b) shows that the preferred nucleation sites of recrystallized grains are at strained grain boundaries, which is a significant feature of the recrystallized structures of IN 718 alloy. The fully recrystallized microstructures were obtained at strain of around 0.7 at temperature of 1050 °C and above, as seen in Fig. 6(c). The fraction of recrystallized grains increased with strain rate in the range of 0.1 to 5 × 10 -3 s -l, but the influence of strain rate on this fraction is much less than the influence of temperature.
4. Discussion 4.1. Flow stress curve behaviour a n d d y n a m i c recrystallization
The flow stress curves of IN 718 alloy developed under high temperature compression conditions, shown in Fig. 2, belong to those of the "normal" be-
Fig. 6. True stress-strain curve and microstructure evolution due to increasing strain, tested at 1050 °C and 0.1 s-~ strain rate: (a) initiation of dynamic recrystallization at the grain boundaries at e = ep; (b) typical necklace structure obtained at ~'= 0.3; (c) fully dynamically recrystallized microstructure achieved at e = 0.75.
haviour of the alloy group with a low or intermediate value of stacking fault energy, and are similar to the curves for nickel [1, 13, 14], copper and austenitic steels [1, 4, 15, 16]. The initial rapid rise in stress is associated with an increase in dislocation density and the formation of poorly developed subgrain boundaries, as a result of work hardening and dynamic recovery. In alloys with a low or intermediate stacking fault energy, e.g. nickel or stainless steel, the dynamic recovery process is very slow, and the dislocation density can attain a sufficiently high value for dynamic recrystallization to be initiated, once a critical strain is exceeded. The role played by the dislocation density in the initiation of dynamic recrystallization has been examined previously [11, 17, 18]. It is apparent that a sufficient energy density must be built up to provide a driving force for the growth of nuclei. It can be seen that ep increased as the strain rate increased and the temperature decreased, as represented in Figs. 2 and 3. The recrystallized grains were observed at ep, as shown in Fig. 6(a) which indicates
6
L.X. Zhou, T. N. Baker
/
Behaviour oflN 718 during high temperature deformation
b the Burgers vector; r the dislocation line energy; dsb the subgrain diameter; m the mobility; and g the strain rate. Hence a critical dislocation density, which depends on the deformation conditions and the grain-boundary characteristics (grain-boundary energy and mobility), must be exceeded before a stable nucleus for dynamic recrystallization can be formed. As a rule, the grain-boundary mobility conforms to an Arrhenius law [22], i.e. m = m 0 exp R T -
Fig. 7. Dynamic recrystallizationinfluenced by the temperature tested at g = 0.1 s-1 and (a) 950 °C, (b) 1050 °C.
that dynamic recrystallization started before peak strain was reached. Under the conditions of strain, nucleation originates at pre-existing grain boundaries, where the dislocation density is sufficient to initiate dynamic recrystallization. The nucleation of recrystallized grains has been found at around 0.5 to 0.7 ep in earlier work [7, 19, 20, 21], and it can be assumed that the critical strain to initiate dynamic recrystallization is equal to •p [19]. It is further indicated that a critical dislocation density, which depends on the conditions of deformation, must be exceeded before recrystallization can proceed. The effects of temperature on the critical strain are similar to those of e_. This can be explained by the analysis of Roberts Iv4], where the dynamic recrystallization occurs at a critical value of dislocation density which is expressed as
p0-const./be,7
) =po
(1>
where P0 is the dislocation density in unrecrystallized material; P0c the critical dislocation density for dynamic recrystallization; s the grain-boundary energy;
(2)
where m0 is a constant. Qm is an activation energy. T and R are, respectively, temperature and gas constant. An increasing temperature leads to an increase in grain-boundary mobility and a decrease in the critical dislocation density for dynamic recrystallization, which in turn results in a decrease in the peak strain ep. Strain rate has the opposite effect. The driving force for the growth of dynamically recrystallized grains is gradient of dislocation density. Figure 5 shows typical features of dynamic recrystallization. A new grain at (R) on Fig. 5(a) grows by consuming adjacent cell structures which contain a much higher dislocation density. In the new grain, dislocations are present, indicating that this microstructure has been developed by dynamic recrystallization. This is because dynamically recrystallized grain growth is accompanied by deformation processing which strains the new grain and increases the dislocations in it. As a result of this current deformation, the dislocation densities in dynamically recrystallized grains are various and change with their extent of strain, as shown in Fig. 5(b). The fraction of dynamically recrystallized grains increases with the temperature, and the fully recrystallized microstructure was obtained at a temperature greater than 1050°C and strain of around 0.7, as shown in Fig. 7. Because the dislocations have increased mobility through cross-slip and climb, this permits attractive interactions to cause annihilation, and a rearrangement into a lower energy, more regular, substructure. Microstructural change with increasing strain is shown in Fig. 6. The progress of dynamic recrystallization is developed by increasing the strain after initiation at ep. The dynamically recrystallized grains are equiaxed with a low dislocation density induced by concurrent deformation. Increasing the strain generates more dislocations in the alloy, which creates a greater driving force for nucleation and grain growth. The absence of 6 phase in the range 1000 to 1100 °C for all strain rate conditions used in this work,
L. X. Zhou, T. N. Baker
/
Behaviour of IN 718 during high temperature de&rmation
illustrates that the solvus temperature of 6 phase is not affected by the strain rate of the starting material with a large grain size and which is 6 phase free. However, for a testing temperature below 6 solvus temperature (1000°C), 6 phase precipitates during deformation, and usually at grain boundaries, as shown in Fig. 4. 4.2. 7he Zener-Hollomon parameter Z The peak stress increases with increasing temperature and decreasing strain rate in Figs. 2 and 3. The data for peak stress, strain to peak stress and Z, obtained for IN 718 alloy, are given in Table 2. The correlation of the peak flow stress ap with temperature T and strain rate ~ during hot working can be expressed by the following mechanical equation of state:
%=AZ"
(3)
or
o,,=Ae"
{,,01
(4)
where A and n are constants. The flow stress increase with strain rate and the effect at constant strain can be expressed by eqn. (5), and the temperature dependence of flow stress can be expressed by eqn. (6).
Equation (4) consists of both following equations:
o=c "
(5)
A(a)=~ee/kr
(6)
where C is a strength constant which depends upon strain, temperature, and alloy composition, and Q is the activation energy for hot working. R is the gas constant, and n' is the strain rate sensitivity of the flow stress. As shown in Fig. 8, eqn. (3) agrees very well with the experimental data. n was determined to be 0.17 _+0.015, which is similar to the data for strain rate sensitivity calculated from the flow strain curves at temperatures 1000 and 1050 °C. The value of Q for IN 718 alloy can be determined from Fig. 8, as Q = 400 +_25 kJ mol-~, in agreement with the results of Guimaraes and Jonas [13] and Camus et al. [14]. The activation energies obtained in hot torsion, hot compression and creep for IN 718 alloy, nickel and nickel base alloys are listed in Table 3. There is a good agreement in the data for Q for IN 718 alloy, and this
103
+ + y
T A B L E 2. Observed values of ap and 0o.7 u n d e r various Z condition
7
+
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70
Strain
Temp.
Z
O'p
a0. 7
rate
(°C)
(s ~:1
(MPa)
(MPa)
1000 1050 950 1000 1050 950 1000 1050
1 . 2 9 4 x 1() 14 3 . 1 0 1 x 1 ( ) Is 6.065 x 1015 1 . 2 9 4 x 1 ( ) ~5 3.101 x 1014 1.213× 1() > 2 . 5 8 8 × 1015 6.203 x 1() 14 1.102 × 1014
152 120 305 200 150 317 225 175 148
112 82 225 148 110 -168 125
(s -~) 5 x 10 5x10 5 x 10 5x10 5 x 10 0.1 0.1 0.1
i ~ : : -"
0.1
1110
100
101 1013
1014
lOX5
1016
1017
1018
Z (l/s)
Fig. 8. Correlation of peak stress with Z e n e r - H o l l o m o n parameter for IN 718 alloy in present work ( ) and previous work conducted by hot compression from ref. 13 (. . . . ) and by hot torsion from ref" 14 ( ..... ).
T A B L E 3. Comparative values of n and activation energy Q Material
Process
Q (kJ tool- 1)
Reference
IN 718 IN 718 IN 718 Ni Ee- 1 8 % C r - 8 % N i Fe- 1 8 % C r - 11%Ni Fe- 1 8 % C r - 1 l%Ni
Hot compression Hot torsion Hot compression Hot working Hot torsion Hot torsion Creep
400 + 25 400 _+ 15 400 + 50 297 414 410 314
Present work [ 14] [ 13] [25] [ 1] [ 15] [ 16]
8
L.X. Zhou, T. N. Baker
/
BehaviouroflN 718duringhigh temperaturedeformation
implies that the activation energy, Q, is not affected by the experimental mode. For nickel base alloys, the values derived from high strain rate tests are much higher than those for creep, because dynamic restoration was by recrystallization in the former and by polygonization in the latter. The value of Q for the pure metal is, as expected, much lower than that for alloys.
4.3. Effect o f strain rate on microstructure
The effect of strain rate in the range from 0.1 to 5 x 10-3 s-1 on the microstructure is not very strong, which is in good agreement with the conclusion reported by Barker et al. [23]. The results are different from those published by Camus et al. in 1986 [14]. There are two reasons for the difference. First, the experiments of Camus et al. were conducted by hot torsion, which is a different mode of testing from the high temperature compression test employed in the present research. Also the flow stress-strain data was translated from the torsion mode to a true stress-strain. In the present work, using the system designed for high temperature compression, the true stress and strain were obtained directly from the recorded data. Comparing the test results from torsion with those obtained from compression, it is found that the torsion data is somewhat higher than the equivalent quantity measured in compression, and similar to that reported previously [24]. Additionally, the range of strain rate in the two studies was different. The work of Camus et al. [14] was carried out using a higher strain rate range from 4 to 4 × 10 -2 s -1, compared with the strain rate range used in the present work which was from 0.1 to 5 x 10- 3 s- i. Also in the work of Camus et al. [ 14], a large amount of metadynamic recrystallization was observed in the specimens produced at the higher strain. Metadynamic recrystallization is a zero incubation recrystallization, based on the existing nuclei which proceed recrystallization and cause an increase in grain growth compared with dynamic recrystallization, because there is no concurrent deformation to restrict grain growth. The strain rate has a very strong effect on the metadynamic recrystallization progress, especially on the delay time (the time from the deformation process ceasing to the metadynamic recrystallization starting); the higher the strain rate, the shorter the delay time. Metadynamic recrystallization could start after less than 1 s at a strain rate as high as 10 s- 1. Because there is some metadynamic recrystallization found in this work, the grain size developed under the higher strain rates is not due exclusively to the dynamic recrystallization is greater than that formed by dynamic recrystallization alone.
o
Conclusions
(1)
In IN 718 alloy, it has been shown experimentally from compression testing in a 200 ton microprocessor controlled press, that flow stresses are temperature and strain rate dependent, and increase with decreasing temperature and increasing strain rate. (2) The relation between peak stress Op and the corresponding regime for dynamic recrystallization can be expressed by the equation Op-~- 0 . 5 Z 0.17
where Z is the Zener-Holloman parameter. (3) The strain rate sensitivities of the flow stress at 1000 and 1050°C are 0.156 and 0.166 respectively. (4) Dynamic recrystallization occurs during the deformation over a temperature range 950 to 1100 °C, and increases with the temperature. A fully dynamically recrystallized microstructure can be achieved at temperatures above 1050 °C. (5) The solvus temperature of 6 phase is not influenced by the strain rates of deformation used in this work. (6) The strain ep to peak stress Op is also affected by the temperature and strain. (7) Initiation of dynamic recrystallization starts at ep and the favourable nucleation sites are the original grain boundaries. (8) Before steady state conditions are reached, increasing strain leads to dynamic recrystallization progress. (9) The rate of metadynamic recrystallization is related to the processing parameters, and increases with temperature and strain rate.
Acknowledgments The authors wish to thank Cameron Forged Products of Livingstone for financial support. Acknowledgments are also made to the Committee of Vice-Chancellors and Principals of the University of the United Kingdom for an ORS award to one of the authors (L.X.Z.).
References 1 J.J. Jonas, C. M. Sellars and W. J. McG. Tegart, Metall. Rev., 14(1969) 1. 2 H.J. McQueen, Microstruct. Sci., 7(1979) 71. 3 C.M. Sellars, Met. Forum, 4(1981) 75.
L. X. Zhou, T. N. Baker
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Behaviour o~ IN 718 during high temperature deformation
4 W. Roberts, in G. Krauss (ed.), Deformation, Processing and Structure, AM. Soc. Metals, Metal Park, OH, 1984, pp. 109-184. 5 T. Sakai and J. J. Jonas, Acre Metall., 32 (1984), 189. 6 M.J. Luton and C. M. Sellars. Acta Metall., 17(196911033. 7 H.J. McQueen, Mater. Sci. Eng. A, 101 (19881 149. 8 T. Sakai, M. Ohashi and K. Chiba, Acta Metall., 36 (19881 1781. 9 L. Blaz, T. Sakai and J. J. Jonas, Met. Sci., 17( 19831609. 10 H.J. McQueen and S. Bergerson, Met. Sci., 6 (1972) 25. 11 J.J. Jonas, l'roc. 4th Int. Coqf Strength of Metals and Alloys, Nan~y, France, 1976, Vol. 2, pp. 976-1001. 12 J. E Collier, A. O. Selius and J. K. Tien, Superalloy 1984, AIME, p. 43. 13 A. A. Guimaraes and J. J. Jonas, Memll. Trans. A, 12 (1981 ) 1655. 14 G. Camus, B. Pieraggi and F. Chevet, Proc. Symp. co-sponsored by the mechanical metallurgy and shaping and forming committees of TMS-AIME, October 1986, pp. 305-325. 15 D.R. Barraclough and C. M. Sellars, Inst. Phys. Conf. Ser., 21, 1974.
9
16 F. Garofalo, C. Richmond, W. F. Domis and F. Von Gemmingen, Joint Int. Conf. ('reep, Institute of Mechanical Engineers, London, 1963, pp. 31-39. 17 R. Sandstrom and R. Langeborg, Acta Metall., 32 (19751 387. 18 W. Roberts and B. Ahlblom, Proc. 4th Int. ('onf. Strength of Metals and Alloys, Nancy, France, 1976, Vol. 1, pp. 400-404. 19 T. Sakai and M. Ohashi, Mater. Sci. Technol., 6 (1990) 1251. 20 J. E Sah, G. J. Richardson and C. M. Sellars, Met. Sci., (19741325. 21 D. W. Livesey and C. M. Sellars, Mater. Sci. Technol., 1 (19851 136. 22 W. Roberts and B. Ahlblom, Acta Metall., 26 ( 197818(11. 23 J.F. Barker, D. D. Krueger and D. R. Chang, Proc. Advanced ttigh Temperature Alloys, ASM, Boston, USA, Oct. 1984, pp. 125-137. 24 J.J. Jonas and T. Sakai, Deformation, in G. Krauss (ed.), Processing and Structure, AM. Soc. Metals, Metal Park, OH, 1984, pp. 185-230. 25 C. M. Sellars and W. J. McG. Tegart, Mere. Sci. Rev, Metall., (~3(1966) 731.