稀有金属(英文版) 2015,34(01),51-63

收稿日期：24 December 2013

基金：financially supported by the Military Supporting Project (No. JPPT125GJGG11);

Hot workability characteristics of Rene88DT superalloy with directionally solidified microstructure

Fu-Lin Li Rui Fu Di Feng Zhi-Ling Tian

Superalloy Department, Central Iron & Steel Research Institute

Abstract：

The hot deformation characteristics of Rene88 DT superalloy with directionally solidified microstructure produced by electroslag remelting continuous directionally solidification(ESR-CDSò) were studied in the temperature range of 1,040–1,140 °C and strain rate range of 0.001–1.000 s-1by hot compression tests. Flow curves for Rene88 DT alloy with initial directionally solidified(DS)microstructure exhibit pronounced peak stresses at the early stage of deformation followed by the occurrence of dynamic softening phenomenon. Rene88 DT alloy with DS microstructure shows higher flow peak stresses compared with HIPed P/M superalloy FGH4096, but the disparities in peak stresses between ESR-CDSed Rene88 DT and HIPed P/M superalloy FGH4096 reduce as temperature increases. The improvement of hot workability of DS alloy with columnar grains avoiding the maximum shear stress comes true. A hot deformation constitutive equation as a function of strain that describes the dependence of flow stress on strain rate and temperature is established. Hot deformation apparent activation energy(Q) varies not only with the strain rate and temperature but also with strain. The strain rate sensitivity exponent(m) map is established at the strain of 0.8, which reveals that global dynamic recrystallization(DRX) shows a relatively high m value in a large strain compression.Optimum parameters are predicted in two regions:T = 1,100–1,130 °C,_e = 0.100–1.000 s-1and T = 1,080–1,100 °C,_e = 0.010–0.100 s-1, which is based on processing maps and deformation microstructure observations.

Keyword：

Rene88DT alloy; Directional solidification; Isothermal compression; Constitutive equation; Processing map;

Author： Fu-Lin Li,e-mail: lifulin1016@sina.com;

Received： 24 December 2013

1 Introduction

Aerospace gas turbine disks operate in an environment ofrelatively high stresses caused by centrifugal forces andelevated temperatures, which necessitate the need formaterials with high-temperature strength and good lowcycle fatigue resistance [1–3]. For example, Rene88DTsuperalloy prepared by power metallurgy (P/M) route iswidely used in the fabrication of turbine disk, owing to itsenhanced balancing and combination of strength, creep,stress rupture, and fatigue crack growth resistance properties [4–6]. Recently, a novel cast and wrought process routefor alloy preparation (registered as ESR-CDS^òalloy) hasbeen designed for the manufacture and applications ofturbine disks [7–9].

In the past few decades, the P/M route was found to beattractive for the manufacture of turbine components in Nibased superalloy [1–5]. Although P/M process overcomesthe severe segregation and produces a fine, homogeneous,and macrosegregation-free structure, the undesirable priorparticle boundaries (PPB) and thermal-induced pores (TIP)cannot be completely eliminated. High alloying degreerestricts the availability of conventional cast and wroughtmetalwork process in the manufacture of turbine disks upto 700 °C and above. An ESR-CDS^òtechnology allowingthe obtainment of high pure and macrosegregation-freebillet containing parallel columnar crystals has beenrecently developed by our research group, which is a newmanufacturing technology combining electroslag remeltingtechnology with directionally solidification technology [7–9]. With continuous withdrawal system, this form ofdirectional solidification (DS) processing was employed sothat heat was withdrawn from the molten metal undercontrol and furthermore promoted the formation ofcolumnar grains. P/M processing of superalloys requiresconsolidation, usually by extrusion or hot isostatic pressing, followed by isothermal forging. However, rapid isothermal forging on the as-cast directionally solidifiedbillets of Rene88DT alloy produced by ESR-CDS^òtechnology could be directly performed under the conditionsof heavy deformation, which significantly lowered themanufacturing costs and cycle compared with P/M route[8, 9]. Thus, there is a critical demand for confirming thedeformation characteristics of ESR-CDSed Rene88DTsuperalloy.

Processing map developed on the basis of dynamicmaterials model by Prasad et al. [10] is a representation ofthe reflection of the material in terms of microstructuralvariation. They were used to optimize the processingparameters and control the microstructure in hot workability of many metallic materials such as P/M superalloy,wrought Ni-base superalloy, aluminum alloys, magnesiumalloys, titanium alloys, and intermetallic alloys [11–20].

As we all know, Ni-base superalloy is rather difficult todeform due to the ever-increasing levels of alloying [6].Workability is dependent on microstructural evolution anddeformation conditions. Furthermore, it is governed by theinitial intrinsic microstructure and the externally imposedstress state, both of which vary with deformation process.Gu et al. [21] developed an advanced cast-and-wrought Ni–Co base superalloy TMW-4M3 for turbine disk applications beyond 700 °C. So far, studies about isothermalcompression behavior for as-cast directionally solidifiedNi-base disk superalloy Rene88DT are rare. Therefore,there is an interest in its as-cast directionally solidifiedsamples, which is the primary coverage of the presentinvestigation. This article intends to give an insight into thedeformation behavior of ESR-CDSed Rene88DT superalloy, and the isothermal compression of ESR-CDSedRene88DT superalloy was conducted at different temperatures and strain rates. The influence of processingparameters on flow behavior, microstructure mechanism,power dissipation (g), deformation activation energy (Q),and strain rate sensitivity exponent (m) was investigated.All of these provide a basis for optimizing the hot workingprocess of ESR-CDSed Rene88DT superalloy.

2 Experimental

In this paper, the Rene88DT superalloy with directionallysolidified microstructure prepared by ESR-CDS^òtechnology was named ESR-CDSed Rene88DT superalloy. Thechemical compositions of the ESR-CDSed Rene88DTalloy used in this study are shown in Table 1. The sourcematerial was duplex-melted (vacuum-induction melting(VIM) and electroslag remelting continuously directionalsolidification (ESR-CDS)). Figure 1 shows the schematicsketch of ESR-CDS^òequipment. During the remeltingprocess, the entire casting was gradually withdrawn fromthe hot zone at a pre-determined rate, with the solid/liquidinterface normal to the withdrawal direction.

In this study, cylindrical billet of Rene88DT alloy with adiameter of 170 mm and a height of 200 mm was cast. Thedirectionally solidified billet macrostructure of ESRCDSed Rene88DT is shown in Fig. 2. As seen from Fig. 2,directionally solidified columnar grains show low orientation deviation angles with billet axial within 15°. Thecolumnar grains at the center are nearly vertical, i.e., normal to the cross section, and the grains at the edge appearto perge away from the center line. This indicates that thesolidification surface is flat at the center and slightly convex at the edges. No evidence of macrosegregation existsin any of the DS casting.

Cylindrical compression samples with a height of15 mm and diameter of 10 mm were cut along with theaxial direction of annealed billets, guaranteeing that thecompression direction for directionally solidified structurewas [001] approximately. The hot compressed tests wereconducted on a Gleeble-3500 thermal simulator at temperatures ranging from 1,040 to 1,160 °C and strain ratesvarying between 0.001 and 1.000 s^-1. Both ends of specimens were lubricated by applying graphite foils. Specimens were heated to the test temperature and soaked for5 min before isothermal compression. The specimens weredeformed to the height reduction of 70 % and then cooleddown to the room temperature immediately. The deformedspecimens were sectioned parallel to the compression axis,and Kallings etchant consisting of 5 g Cu Cl₂, 100 ml HCland 100 ml ethanol was used for revealing the microstructure. Dynamic recrystallization microstructural examination was conducted by Olympus-PMG3 opticalmicroscope (OM). The microstructures of as-cast and asannealed billet were characterized by EVO 18 scanningelectron microscopy (SEM).

Fig. 1 Schematic sketch of equipment of ESR-CDS?

Table 1 Nominal compositions of ESR-CDSed Rene88DT (wt%)  下载原图

Table 1 Nominal compositions of ESR-CDSed Rene88DT (wt%)

Fig. 2 Directionally solidified billet macrostructure of ESR-CDSedRene88DT

3 Results and discussion

3.1 Microstructures of DS billet

The as-cast microstructure of ESR-CDSed Rene88DTsuperalloy shown in Fig. 3a exhibits a typical dendritestructure consisting of dendrite region marked as A andinter-dendrite region marked as B. Its columnar grainedstructure with [001] direction parallels to the axial directionof solidification billet. During the continuous DS process, ahigher thermal gradient and solidification velocity canresult in smaller dendrite spacing, and the secondary dendrite arm spacing (C) ranges from 50 to 100 lm. Thedecrease of secondary dendrite spacing makes it possible toreduce the segregation degree of billet microstructure so asto obtain the high quality alloy casting ingot. As shown inFig. 3b, diffusion that occurs during the solution annealingheat treatment at 1,200 °C for 24 h reduces the segregationdegree between dendrite arms and interdendritic regionsand results in more homogeneous element distribution.

Fig. 3 SEM images of ESR-CDSed Rene88DT alloy: a columnargrained structure of longitudinal section of billet before annealing and

Fig. 3 SEM images of ESR-CDSed Rene88DT alloy: a columnargrained structure of longitudinal section of billet before annealing andb homogenous microstructure after annealing at 1,200 °C for 24 h

3.2 Flow behavior

The typical true stress–true strain curves of the ESRCDSed Rene88DT alloy obtained at different deformationtemperatures from 1,040 to 1,140 °C with various strainrates from 0.001 to 1.000 s^-1are shown in Fig. 4. As seenfrom Fig. 4, work hardening is notable at the initial stage ofthe deformation, which leads to a pronounced peak stress.As shown in Fig. 4a–d, at high strain rates, the flow curvesexhibit strain hardening to a peak followed by work softening to a steady-state regime. At low strain rates, anextensive steady stress state is quickly achieved in whichthe flow stress almost keeps unchanged with the increase ofstrain. The features of the flow curves for ESR-CDSedRene88DT alloy indicate the occurrence of DRX phenomenon. DRX is shown to be a more efficient process tokeep the flow stresses and the rates of work hardeningconsiderably low.

It can be noted that, different from the other temperatures,at 1,120 and 1,140 °C, flow curves of the deformationexhibit slow hardening trends or bulk steady states, whichcan be explained by the relation of interaction betweendislocation motion and c⁰phase [22]. The amount of c⁰phase decreases with temperature increasing. When thedeformation temperature is below 1,110 °C which is thecomplete dissolution temperature of c⁰phase, undissolved c⁰phases hinder the dislocation motion at the early stage ofcompression. Once a critical strain is reached, the proliferation, the annihilating, and the mutual reactions of the dislocations stimulate the occurrence of DRX. It is clear that c⁰phase remarkably affects the DRX, so that distinct difference exists in the flow behaviors between above and below1,110 °C. Figure 5 shows the variation in the peak stress asa function of temperature at different strain rates. The peakstress increases with strain rate and decreases with temperature. The influence of the strain rate on the peak stress ismore pronounced at low temperatures, and the difference inthe peak stress between various temperatures reduces withstrain rates decreasing. In fact, higher temperature and lowerstrain rate can make energy accumulation and mobility atboundaries for the nucleation and growth of dynamicallyrecrystallized grains easier, thus reducing the flow stresslevel. It is concluded that the flow behaviors of alloy areessentially dominated by dynamic recovery, dynamicrecrystallization, and dislocation mechanism.

Fig. 4 True stress–true strain curves of ESR-CDSed Rene88DT alloy at various temperatures: a 1,040 °C, b 1,060 °C, c 1,080 °C, d 1,100 °C,e 1,120 °C, and f 1,140 °C

The peak stress in the temperature range of 1,060–1,140 °Cand strain rate range of 0.001–1.000 s^-1for ESR-CDSedRene88DT superalloy and hot isostatic pressed P/M superalloy FGH4096 is given in Fig. 6. FGH4096, type in China,possesses the same chemical composition as Rene88DTsuperalloy [23–25]. It can be noted that with the temperatureincreasing, disparities in the peak stresses between ESRCDSed Rene88DT and HIPed P/M superalloy FGH4096 atthe same strain rate reduce. At higher deformationtemperatures (1,120–1,140 °C), the peak stress values of thetwo alloys with different initial microstructures are verysimilar. Because of original larger DS columnar grains, thecompression curves of ESR-CDSed Rene88DT show higherpeak stresses compared with those of HIPed P/M superalloyFGH4096. For as-cast directionally solidified alloyRene88DT, there exist paralleling large grains with the size ofseveral millimeters level. This determines that the onset ofdynamic recrystallization in as-cast directionally solidifiedalloy Rene88DT is nucleation-controlled instead of growthcontrol for polycrystalline materials. When the DRX grainsare larger, the main deformation mechanisms are dislocationaccommodated grain boundary sliding and dislocation slipplasticity. At higher deformation temperature, the thermalactivation can rapidly stimulate the motion of dislocations andthe peak stress is reached soon followed by the occurrence ofDRX. In the process of axial compression, these directionallysolidified columnar grains without transverse grain boundaries can avoid the maximum shear stress direction, i.e., thedirection of 45° with principal stress axial direction. Consequently, the hot working plasticity of ESR-CDSed Rene88DTalloy can be effectively improved.

Fig. 5 Peak stress as a function of temperature at different strain rates

Fig. 6 Peak stress under different deformation temperatures andstrain rates for superalloy ESR-CDSed Rene88DT and hot isostaticpressed FGH4096

3.3 Development of constitutive equation

The relationship between flow stress, strain rate, anddeformation temperature can be described by a hyperbolicsine Arrhenius-type equation proposed by Sellars and Tegart [26]:

where A (s^-1) and a (MPa^-1) are material constants, n is aconstant related to processing parameters, Q (J mol-1) isthe activation energy of deformation, R is the universal gasconstant, T is the deformation, r (MPa) is the flow stress,and e_ (s^-1) is the strain rate.

As is known, Eq. (1) approximates a power law Eq. (2)when ar \ 0.8 and an exponential law Eq. (3) whenar [ 1.6:

where b and n₁are also constant related to processingparameters. Taking the logarithm of both sides of Eqs. (2)and (3), we can get Eqs. (5) and (6):

Under the corresponding true strain ranging from 0.2 to1.2, the values of n₁can be determined by Eqs. (5) and (6)from the slopes of the lines of r - lne _ and ln r - lne _.Figure 7 shows the regression analysis data and results atthe strain of 0.5 under various deformation temperaturesand strain rates, which show good linear relationships.Then the value of a can be calculated by Eq. (4). It isconcluded that the value of a remains unchanged at eachcorresponding strain ranging from 0.1 to 1.2, and a value isapproximately estimated to be 0.0063 MPa^-1.

Taking the natural logarithms of both sides of Eq. (1),yields the following:

The plot of ln e_ versus lnsinh(ar) and the plot oflnsinh(ar) versus 1,000 T^-1at peak strains are shown inFig. 8a and b, respectively. The value of Q can beestimated by the slope of the plot of lnsinh(ar) versus1000 T^-1at a constant strain and strain rate. Thecalculation of Q from Eq. (8) can be obtained bydifferentiating Eq. (7).

At temperatures where thermally activated deformationand restoration processes occur, the microstructuralevolution will be dependent on the deformationtemperature (T) and strain rate ( e_) in addition to thestrain (e). The strain rate and deformation temperature areoften incorporated into a single parameter—the Zener–Hollomon parameter (Z), which is defined as

Fig. 7 Plots of r versus ln e_ a and lnr versus lne _ b at e = 0.5 under different temperatures

Fig. 8 Relationships of hot deformation peak stress with strain rates a and temperatures b of ESR-CDSed Rene88DT alloy

Fig. 9 Plot of flow stress as a function of Zener–Hollomon parameterat strain of 0.2

Parameter Z can be used to demonstrate the combinedeffect of strain rate and temperature on the hot deformationbehavior of alloy material. Figure 9 shows the relation offlow stress with the parameter Z for the ESR-CDSedRene88DT alloy at the strain of 0.2, which exhibits goodlinear relationship with a correlation coefficient of 0.9839.The stress increases consistently with the Zener–Hollomonparameter increasing for the given test conditions.

3.4 Strain-dependent constitutive analysis

Since the flow stress may be strain-dependent during hotdeformation, calculating that the flow stress using constitutive equation with stain involved is inevitable [27]. Theparallel lines obtained from the plots of ln e_ versus lnsinh(ar) at various strains using a single a value can justify avalue as constant. In this study, the calculation resultsreveal that a value can be determined to be constant in thewhole deformation process. Figure 10 shows the values ofmaterial constant (Q, ln A and n) under different deformation strains within the range of 0.2–1.0 and with theinterval of 0.1. At low strain regions, the n, ln A, andQ values decrease with the strain increasing. At highstrains, Q and ln A curves show slow decreasing trends,whereas the n value increases with the strain increasing. Itcan be interpreted that at a given temperature and strainrate, the increase of n value has a compensation effect onthe decline of flow stress related to dynamic softening. Thevariations in n value are not significantly sensitive to thestrain, but they cannot be treated as constants. If they areseen as constants, the variations in flow stress at differentstrains will follow the changes in A value under all testconditions. As hot deformation is a thermally activatedprocess, accompanied with grain rotation, more slip systems are activated when the strain increases. It results inthe decrease of activation energy, making it easier for theprogress of hot deformation. Consequently, the constitutiveparameters are strain-dependent. The relationships betweenQ, ln A, n, and true strain can be polynomial fitted by thecompensation of strain in this study. The strain-dependentconstitutive equation for ESR-CDSed Rene88DT alloy canbe written as follows:

Fig. 10 Variations in constitutive parameters as a function of strain rate: a Q, b n, and c lnA

Fig. 11 Effect of strain on activation energy at a strain rate of 0.100 s-1and deformation temperatures of 1040, 1100, and 1140 °C andb deformation temperature of 1,080 °C and strain rates of 0.001, 0.010, 0.100, and 1.000 s-1

3.5 Variations of activation energy with deformationparameters

In the current study, it is found that deformation activationenergy was significantly influenced by strain rate, deformationtemperature, and strain. The dislocation motion, dynamicrecovery (DRV), and DRX were related to activation energy,which reflected the deformation behaviors of alloy materials.The effect of strain on activation energy at various temperatures and strain rates is shown in Fig. 11. Nucleation of crystallographic defects such as dislocations required high thermalactivation energy at deformation incipient period. Meanwhile,the interactions of dislocations with other dislocations, grainboundaries, c⁰phases, and periodic friction of lattice itself canresult in a small increasing trend of activation energy in theinitial stage of strain as shown in Fig. 11. During the compressive deformation, more slip systems are gradually activated. Once the energy barrier is overcome with the assistanceof a combination of both thermal and mechanical stresses,activation energy shows a continuous declining tendency withthe further increase of strain, and DRV and DRX occur.

Figure 12 shows the activation energy–strain rate curvesat strains of 0.2, 0.5, 0.8, and 1.0 and deformation temperatures of 1040, 1080, 1100, and 1140 °C. It can be seenthat the activation energy exhibits relatively low valueswithin the strain rate range of 0.010–0.100 s^-1at strains of0.5, 0.8, and 1.0 and various temperatures. The valleyvalues of activation energy are attributed to the mechanismof DRX and DRV that correlate with the decrease indensity of dislocation and substructure. At the lowest strainof 0.2, however, activation energy is insensitive to strainrate at the given deformation temperatures and strain ratesfrom 0.001 to 1.000 s^-1. At all test conditions, the activation energy increases with the increase in strain ratewhen strain rate is higher than 0.100 s^-1, which can beexplained by the increase of dislocation density and average velocity of dislocations moving though a field ofobstacles. Therefore, to decrease the activation energy, thedeformation optimum strain rate can be intended to be inthe range of 0.010–0.100 s^-1.

Fig. 12 Effect of strain rate on activation energy at strains of 0.2, 0.5, 0.8, and 1.0 and deformation temperatures of a 1,040 °C, b 1,080 °C,c 1,100 °C, and d 1,140 °C

With the above analysis, we also present the influenceof deformation temperature on activation energy at astrain rate of 0.100 s^-1and strains of 0.2, 0.5, 0.7, and 1.0as shown in Fig. 13. As seen from the activation energycurves, the valley values emerge in the vicinity of temperature range of 1,080^-1,100 °C. Overall, the effect ofdeformation temperature on activation energy is similar atvarious strains with the strain rate of 0.100 s^-1. Under thecritical temperature of c0phase full dissolution, a certainamount of c0phase can significantly stimulate the nucleation of recrystallized grains and hinder their growthconcurrently, which can maintain a lower average dislocation density. For the compressive deformation at alower temperature of 1,040 °C, slip systems havedifficulty in activation because of the insufficient thermalenergy. Yet at the elevated temperature of 1,140 °C, as allc⁰phases are dissolved, recrystallized grain coarsening isapparently promoted. Grain coarsening can cause workhardening, and boundary migration will require higheractivation energy. In order to decrease the activationenergy and achieve the well-refined recrystallized grains,deformation temperature region should be in the range of1,080–1,100 °C.

Fig. 13 Effect of deformation temperature on activation energy atstrain rate of 0.100 s-1and strains of 0.2, 0.5, 0.7, and 1.0

3.6 Processing map

Briefly, the workpiece undergoing hot deformation isconsidered to be a dissipater of power supplied by a particular source. The power (P) dissipated by the workpiececould be partitioned into two complementary parts calledpower content (G) and power co-content (J).

The first integral is defined as G content and the secondis defined as J co-content, which is related to the mainpower input dissipated in the form of a temperature rise andthepowerdissipatedbymetallurgicalprocesses,respectively. At a constant strain and temperature, flowstress could be represented by Eq. (12), and strain ratesensitivity (m) can be obtained by Eq. (13).

where K is a material constant.

According to Eqs. (11) and (12), the J content isevaluated as

So the maximum value of J could be represented as

The efficiency of power dissipation can be derived by

The variation of g with deformation temperature andstrain rate constitutes the power dissipation map. Based onthe extremum principles of irreversible thermodynamic asapplied to large plastic flow, the instability criterion nee _T isgiven by

where flow instabilities are predicted to occur when ne e_Tbecomes negative. The variation of instability criterionparameter (n) with deformation temperature and strain rateconstitutes a flow instability map. The processing map canbe obtained by superimposition of power dissipation efficiency map and flow instability map.

Using the cubic spline interpolation, the flow stress at finerintervals of temperature and strain rate were calculated. Atevery sub-interval of temperature, the strain rate sensitivityparameters (m) were calculated using Eq. (13). The variationof strain rate sensitivity (m) with respect to temperature andstrain rate is given in Fig. 14. It can be observed that over theentire domain, m varied from 0.13 to 0.32. The values ofm increase with the temperature increasing from 1,040 to1,100 °C and strain rates decreasing.

Processing map is an explicit representation of theresponse of the material in terms of microstructural mechanism to the imposed process parameters. The processing mapcan provide domains where a local maximum occurs in theefficiency of power dissipation and these are the zones wherethe dissipative energy of the material is the lowest. Furthermore, since the energy conversion into surface energy is mostefficient, the efficiency of power dissipation by crackingprocess is also generally very high. In brief, the processingmaps consist of a superimposition of power dissipation mapand instability map, which can illustrate the‘‘safe’’domainand‘‘unsafe’’domain during plastic processing. Power dissipation occurs by complementary processing: a large part asheat through plastic deformation and the other part throughmicrostructural changes. Previous study revealed that thestrain had no significant role in the processing maps.

Based on the above dynamic material model and experimental results of isothermal compression of ESR-CDSedRene88DT alloy, the map of power dissipation efficiency at atrue strain of 0.8 can be obtained as shown in Fig. 15. Figure 15can also be thought as power dissipation maps and can beinterpreted in terms of the microstructural processes. Generally, the map exhibits four domains in the temperature andstrain rate ranges: I, T = 1,040–1,060 °C and e_ = 0.001–1.000 s^-1with lower efficiency; II, T = 1,060–1,160 °C and

Fig. 14 Strain rate sensitivity contour (m) maps of ESR-CDSedRene88DT alloy at strain of 0.8

Fig. 15 Power dissipation map of ESR-CDSed Rene88DT alloy atstrain level of 0.8. Numbers representing efficiency of powerdissipation

Fig. 16 Instability criterion parameter map of ESR-CDSedRene88DT alloy at strain level of 0.8

e_ = 0.001–0.030 s^-1with higher efficiency; III, T = 1,140–1,160 °C and e_ = 0.030–1.000 s^-1with lower efficiency; IV,T = 1,080–1,130 °C and e_ = 0.100–1.000 s^-1with higherefficiency. Figure 16 displays the instability map at the strainlevel of 0.8, which reveals that no instability region is foundwithin the selected deformation conditions of this study. Wuet al. [28] studied the hot workability of a new Ni–Co–Cr-basedP/M superalloy and concluded that the samples had obviouslongitudinal cracks when deformation strain value exceeded0.5. PPB and large and hard inclusions limited the deformationstrain. The nickel-base disk alloy with pure and low segregation microstructure produced by ESR-CDS technology showssuperior hot workability. Figure 17 shows macrographs of thecompression samples under various test conditions. It revealsthat all samples exhibit better plasticity even for samplesdeformed in Domains I and III. According to Figs. 14, 15, and16, Domains I and III have low values of m, g, and ne e_T.Therefore, a feasible processing window deduced fromFigs. 14, 15, and 16 consists of two regions: T = 1,100–1,130 °C and e_ = 0.100–1.000 s^-1, and T = 1,080–1,100 °Cand e_ = 0.001–0.100 s^-1. To investigate the microscopicdeformation mechanisms and verify the reliability of processparameters predicted by processing map, the evidence ofdeformation in these domains is identified and validatedthrough microstructure observations in the following sections.

3.7 Microstructure analysis

The microstructural processes can interpret the domains ofthe power dissipation maps based on the characteristicefficiency variation. The deformation mechanisms of stableregions include dynamic recovery, DRX, and superplasticity. These are desirable processes during hot working ofa uniform billet. DRX can keep flow stresses and rates ofwork hardening considerably low and is considered as apreferred choice for hot working. Generally, the efficiencyvalues associated with DRX are about 30 %–50 %, whilethe values related to superplasticity are usually more than60 %. The OM images of the studied alloy deformed to thestrain level of 1.2 are shown in Fig. 18. The changes in themicrostructure can show evidence concerning the specifiedmechanism dominating the domains of the power dissipation maps.

The microstructural manifestation of processing mapwith lower instability value regarding the present alloy isshown in Fig. 18a, which corresponds to the specimendeformed at 1,060 °C and 0.010 s^-1. The microstructureexhibits bands of flow localization associated with slipsystems which do not operate at low temperatures becauseof a higher Peierls–Nabarro stress. The microstructure ofthe studied alloy deformed at 1,060 °C and 0.010 s^-1(Fig. 18a) corresponding to Domain I of Fig. 15 shows anincomplete recrystallized structure. For Domain I, themicrostructure consists of the necklace structures with finegrains with g value of less than 30 %. It suggests that thisregion mainly represents partial DRX. Uniform and finerecrystallized microstructures with corrugated grainboundaries can be identified in Fig. 18b–d, which arelocated in Domains II and IV of Fig. 15. It implies that theworkability of as-cast ESR-CDS Rene88DT alloy will beimproved in this region. It is widely recognized that highpeak power dissipation efficiency is often associated withdynamic recrystallization or superplasticity. It reveals thatthe two domains, with higher efficiency, can be interpretedto represent full DRX. Moreover, the strain rate sensitivity(m) at the two regions shows higher value as shown inFig. 14. Deformation at the temperature of 1,080–1,110 °Cand strain rate of 0.010–0.001 s^-1shows the peak value ofm. Krueger [5] concluded that certain alloys with a strainrate sensitivity value (m) at preselected working conditions,of at least about 0.3 for a given strain rate, would not resultin critical and abnormal grain growth. However, owing tothe high temperature of 1,160 °C as well as the low strainrate of 0.010 s^-1, the c⁰phase seems to be almost completely dissolved and the grain growth is apparently promoted as shown in Fig. 18f. This kind of coarsening grainmicrostructure corresponds to Domain IV with the lowestvalues of strain rate sensitivity (m) and instability criterionparameter (n).

Fig. 17 Macrographs of isothermal compression samples of ESR-CDSed Rene88DT alloy deformed to height reductions of 70 % under varioustest conditions: a T = 1,040 °C, e_ = 1.000 s^-1; b T = 1,080 °C, e_ = 0.100 s-1; c T = 1,100 °C, e_ = 0.100 s^-1; d T = 1,160 °C,e_ = 0.010 s^-1

Fig. 18 OM images of studied alloy deformed to strain of 1.2 at different deforming conditions: a T = 1,060 °C, e_ = 0.010 s^-1; bT = 1,080 °C, e_ = 0.001 s^-1; c T = 1,100 °C, e_ = 0.001 s^-1; d T = 1,120 °C, e_ = 0.100 s^-1; e T = 1,140 °C, e_ = 0.010 s^-1; f T = 1,160 °C,e_ = 0.010 s^-1

Figure 19 shows the typical microstructure of thespecimen deformed under the temperature of 1,100 °C andthe strain rates of 0.100, 0.050, 0.020, and 0.001 s^-1.Microstructure observation indicates that the DRX grainsizes decrease with strain rate increasing at the temperatureof 1,100 °C. Under higher strain rates, the dislocations arerapidly stimulated and subgrain rotations occur, whichcontributes to the promoted nucleation of DRX [29]. Thereis no enough time for atom diffusion and boundarymigrations, thus the grain growth can be inhibited at strainrates higher than 0.050 s^-1. In combination with the powerdissipation map, it can be seen that Domain II of Fig. 15with the efficiency of higher than 0.4 corresponds to theDRX grain growth.

Fig. 19 OM images of studied alloy deformed to true strain of 1.2 at temperature of 1,100 °C and various strain rates: a 0.100 s-1, b 0.050 s-1,c 0.020 s-1, and d 0.001 s-1

In view of the analysis above, the microstructures by andlarge correspond to the regions of hot processing maps withinthe experimental conditions. Based on the comprehensiveconsideration in the effect of strain rate, temperature and strainon flow stress, activation energy (Q), strain rate sensitivity(m), power dissipation (g) map, instability map, and microstructure observations, hot deformation can be carried outunder conditions of T = 1,100–1,130 °C, e_ = 0.100–1.000 s^-1and T = 1,080–1,100 °C, e_ = 0.010–0.100 s^-1toobtain desirable microstructure.

4 Conclusion

The isothermal compression of ESR-CDSed Rene88DTsuperalloy was studied via physical experiment and microstructure characterization. Flow curves of isothermallycompressed ESR-CDSed Rene88DT superalloy exhibit thecharacteristics of initial work hardening and then dynamicsoftening. Rene88DT alloy with DS microstructure showshigher peak stress compared with HIPed P/M superalloyFGH4096, but the disparities in peak stresses between ESRCDSed Rene88DT and HIPed P/M superalloy FGH4096reduce as temperature increases. The improvement of hotworkability of directionally solidified alloy with columnargrains avoiding the maximum shear stress can be reached.

Based on the experimental stress–strain data, a constitutiveequation is established. In this study, the stress multiplier (a) inthe hyperbolic sine equation is believed to be constant, and theother parameters are made to be strain-dependent. The strainrate, deformation temperature, and strain have profoundeffects on activation energy (Q). The value of Q decreaseswith the increase in strain. No instability zone is detectedunder the selected deformation conditions in the present study.The power dissipation map at the strain level of 0.8 is pidedinto four domains. Optimum parameters are predicted in tworegions: T = 1,100–1,130 °C, e_ = 0.100–1.000 s^-1, andT = 1,080–1,100 °C, e_ = 0.010–0.100 s^-1, which is basedon processing maps and deformation microstructureobservations.