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Evolution of precipitates in Ni-Co-Cr-W-Mo superalloys with different tungsten contents

Han-Sheng Bao Zhi-Hua Gong Zheng-Zong Chen Gang Yang

Institute for Special Steels,Central Iron and Steel Research Institute

School of Materials and Metallurgy,Inner Mongolia University of Science and Technology

作者简介：*Han-Sheng Bao,e-mail:Baohansheng@nercast.com;

收稿日期：6 November 2019

基金：financially supported by the National Key Research and Develop Program,China(No. 2017YFB0305203);

Evolution of precipitates in Ni-Co-Cr-W-Mo superalloys with different tungsten contents

Han-Sheng Bao Zhi-Hua Gong Zheng-Zong Chen Gang Yang

Institute for Special Steels,Central Iron and Steel Research Institute

School of Materials and Metallurgy,Inner Mongolia University of Science and Technology

Abstract：

The Ni-Co-Cr-W-Mo system is critical for the design of nickel-based superalloys.This system stabilizes different topologically close-packed(TCP) phases in many commercially superalloys with high W and Mo contents.Scanning electron microscopy(SEM),transmission electron microscopy(TEM) and thermodynamic calculations were applied to investigate the thermodynamics of the precipitates in two different W-contained Ni-Co-Cr-WMo superalloys(Alloy 1 and Alloy 2).Computational thermodynamics verifies the experimental observation of the μ phase formation as a function of temperature and alloy chemistry,but the kinetics for the precipitation of the M6 C phase do not agree with the experimental findings.The major precipitates of Alloy 1 at temperatures of700℃ and 750℃ during long-time exposure are M23 C6,γ′ phase and MC;for Alloy 2,they are M23 C6,γ′ phase,MC,M6 C and μ phase.W addition is found to promote the precipitation of M6 C and μ phase during exposure.M6 C has higher W and lower Ni content than μ phase,whereas M6 C is an unstable phase that would transform into M12 C after 5000-h exposure at 750℃.A great quantity of needle-like μ phases precipitated after exposure at 750℃ for5000 h,which have no effect on the impact properties of Alloy 2.

Keyword：

Nickel-based superalloys; M₆C; μphase; Long-time exposure; M₁₂C;

Received： 6 November 2019

1 Introduction

With the development of advanced ultra-supercritical coalfired power plants,the steam pressure and temperature reach 35 MPa and 700℃,respectively,while thermal efficiency will be over 50% [ 1, 2, 3, 4, 5] .Conventional Fe-based heat-resistant steels used for main steam piping,and reheater (or super-heater) tubes and blades do not comply with the manufacturing requirements of key high temperature components,for steam temperature exceeding 650℃ [ 6, 7] .Thus,some nickel-based alloys including Waspaloy,USC 141,ACC 617 and Inconel 718 are candidates for turbines and pipes [ 8, 9, 10] .

Tungsten (W) has a larger atomic radius and lattice distortion than molybdenum (Mo) to produce a stronger solution strengthening effect.Additionally,the reduced stacking fault energy of alloys containing tungsten can improve creep fracture strength [ 9, 11] .As a result,W has partially replaced Mo in some heat-resistant steels,such as the TOS110 and HR 1200 [ 12] ,to improve creep rupture strength.In addition,W is also widely used in nickel-based single crystal alloys to increase the service temperature [ 13, 14, 15] .

Waspaloy is a nickel-based superalloy strengthened by nanometer-sizeγ'phase and M₂₃C₆ carbide.As W can improve the stability ofγ'phase through changing the distribution of the alloy elements inγ'-γmatrix,some W element is added to the Waspaloy nickel-based superalloy.Our group recently developed a novel wrought nickelbased superalloy based on Waspaloy that is a promising candidate for use in steam turbine.The addition of excess refractory elements such as W and Mo to the alloy will promote the formation of a phase family known as topologically close-packed (TCP) phases during long-time exposure at high temperature [ 16, 17] .Belan [ 18] reported that TCP phases are primarily plate or needle-like and will reduce rupture strength and ductility.

This work investigated the microstructural evolution in nickel-based alloys with W content varying during longterm exposure at 700-750℃and the influence on the mechanical properties of the alloys.The structure and composition of the various alloys’phases were determined,and the role of W was discussed in terms of the formation of precipitates present in the alloys.

2 Experimental

The main alloying addition of these two alloys including nickel (Ni),cobalt (Co),W,Mo and chromium (Cr) were molten using a vacuum induction furnace.The 25-kg ingots were then forged into a round bar with 16 mm in diameter.The chemical compositions of the new designed alloys are given in Table 1.

The as-processed samples were subjected to a standard heat treatment process,which consists of (1) soaking at1080℃for 4 h and oil cooling,(2) soaking at 845℃for24 h and then air cooling to room temperature and (3)aging at 760℃for 16 h and then air cooling to room temperature.The as-treated samples were subsequently aged in a muffle furnace.Both alloys were aged at 700℃for 1000,3000,5000 and 10,000 h and 750℃for 3000and 5000 h.

Charpy impact properties were evaluated for the samples subjected to standard heat treatment and long-term exposure.The absorbed energies were obtained using a Charpy V-notch machine at room temperature.

Precipitates extracted from aged samples were identified by a small-angle X-ray diffraction (XRD,Pert MPD) tester.Precipitates powder for XRD tests was prepared with a solution consisting of 10 g·L^-1 ammonium sulfate and10 g·L^-1 aqueous citric acid,and the extraction current density was 0.03 A·cm² at a temperature range from-15℃to 20℃.Specimens for field emission scanning electron microscopy (FESEM) were cut from aged samples and prepared using mechanical polishing and etching with a solution containing 200 ml hydrochloric acid,200 ml ethyl alcohol and 10 g copper dichloride.SEM tests were carried out using a FEI Quanta 650 microscope equipped with an energy-dispersive spectrometry (EDS).Thin films for transmission electron microscopy (TEM,JEM-21 00F)specimens were prepared using mechanical milling and conventional dual-jet electro-polishing in a solution of HClO₄+CH₃CH₂OH at a temperature of-30℃and then further thinned by ion beam milling.

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Table 1 Chemical compositions of alloys (wt%)

3 Results and discussion

3.1 Thermo-calc dynamics calculation

Thermo-calc dynamic calculation software predicts the phase formation and stability of alloys [ 19, 20] .In this study,thermal-calc thermodynamic software with TTNi 8database was used to calculate the content of precipitation of alloys with different W and Mo contents at varying temperatures.Figures 1 and 2 show the calculated results of the tentative alloys.Figure 1a shows all the predicted equilibrium phases and their content in Alloy 1,which contains 4 wt%W at a temperature range of 500-1500℃.Only the liquid phase forms at 1380℃(above the liquidus).Theγmatrix,μphase,γ'phase and M₂₃C₆ carbide have been separated out in turn with temperature decreasing and their contents change with temperature.The calculated precipitation temperature ofμphase,γ'phase and M₂₃C₆ phase is 1130,1010 and 920℃,respectively.Figure lb shows all the predicted equilibrium phases and their contents in Alloy 2,which contains 8 wt%W.One can see that there are y matrix,μphase,γ'phase,a phase,R phase and M₂₃C₆ carbide.And the precipitation temperature ofσphase and R phase is 725℃and 580℃,respectively.Meanwhile,a very little MC phase precipitates from two alloys at temperature above 900℃.

Figure 2a,b shows the effect of W content on the equilibrium precipitates of the alloy with 3.5 wt%Mo addition at 700℃and 750℃,respectively.At 700℃,the precipitation amount ofμincreases along with W content when the latter is less than 6 wt%,and then remains unchanged.At 750℃,the precipitation amount ofμincreases with W content,while the latter is less than 8.2wt%,and then remains constant.Meanwhile,with the same W content,the precipitation amount ofμphase at 700℃is higher than that at 750℃.SIGMA phase begins to precipitate when W content reaches 7.8 wt%and 8.3 wt%at700℃and 750℃,respectively.

Figure 3 presents XRD results of the alloys with different W contents after long-term exposure at 700℃and750℃.Onlyγ'phase,MC and M₂₃C₆ carbides were detected in Alloy 1 after exposure at 700℃and 750℃,whereasγ'phase,MC,M₂₃C₆ and M₆C were detected in Alloy 2 before and after long-term exposure;μphase was found after exposure at 750℃for 3000 and 5000 h.SIGMA and R phases were not detected in the present work based on the chemical phase analysis method.

Fig.1 Calculated phase content as a function of temperature diagrams for a Alloy 1 and b Alloy 2

Fig.2 Calculated phase content as a function of W content at a 700℃and b 750℃

Fig.3 XRD patterns of precipitates

Computational thermodynamics verified experimental observations ofγ',MC,M₂₃C₆ andμphase formation,but the amount of MC is minute and does not appear in some calculation results.M₆C is an unstable phase that is possibly a mesophase ofμphase [ 21] .As such,it cannot appear in the equilibrium phase calculation results.

3.2 Microstructural evolution during exposure

Figure 4 shows SEM images of the two alloys after expose for 0,5000 and 10,000 h at 700℃.Ti-rich MC particles were distributed randomly inner on the grains of both Alloy1 and Alloy 2 and maintained stability during exposure as marked by the arrows.Most Cr-rich M₂₃C₆ particles were discontinuously distributed on grain boundaries before exposure and then transformed to have a continuous distribution after exposure for 10,000 h.XRD identified that the W-rich microscale round white particles in Alloy 2 as M₆C arranged in array on the grain as marked by the arrows.

Figure 5 shows SEM images of the alloy after exposure for 3000 and 5000 h at 750℃.The size,amount and morphology of MC carbides do no undergo remarkable changes after 5000-h exposure in either alloy,but the amount of M₂₃C₆ carbides on the grain boundary increased significantly with exposure time increasing and caused grain boundaries to widen.Small needle-like precipitates were observed in Alloy 2 after exposure for 3000 h,and the amount of needle-like precipitation increases significantly with longer exposure time,as shown in Fig.5d.XRD analysis shows that the needle-like precipitations areμphase,randomly distributed in grains.

Fig.4 SEM images showing microstructural evolution of alloys after exposure at 700℃for different periods:Alloy 1 for a 0 h,b 5000 h and c 10,000 h;Alloy 2 for d 0 h,e 5000 and f 10,000 h

Fig.5 SEM images of alloys after exposure at 750℃for different periods:Alloy 1 for a 3000 h and b 5000 h;Alloy 2 for c 3000 and d 5000 h

3.3 Evolution of M6C

When W is high enough,M₆C precipitates at high temperature rather than M₂₃C₆ [ 22, 23] .Figure 6 shows the precipitation evolution during standard heat treatment of Alloy 2.The W-rich white particles precipitate on the grain after solution treatment at 1080℃for 4 h.The amount of M6C increases slightly after pre-exposure at 845℃for24 h,and the M₂₃C₆ carbides begin to precipitate on grain boundaries.The morphology and quantity of M₆C carbides remain unchanged during aging at 760℃for 16 h.Table 2shows the chemical compositions of M₆C and M₂₃C₆obtained by EDS.The total elemental contents of W and Mo of Positions A,B and C in Fig.6 are 56.45 wt%,77.98wt%and 52.24 wt%,respectively.Position D contains15.75 wt%W and Mo and contains more Cr and Ni than the other particles.By comparing the elemental composition of Positions A,B,C and D and combining their XRD results,Positions A,B and C are M₆C and Position D is M₂₃C₆.

Fig.6 SEM images of M₆C and M₂₃C₆ after different heat treatments:a 1080℃,b 1080℃+845℃,and 1080℃+845℃+760℃

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Table 2 Elemental component of different positions in Fig.6 with SEM+EDS analysis (wt%)

Research has shown that M₆C is an unstable phase that is transformed into M₁₂C during long-term exposure at high temperatures [ 24] .M₆C and M₁₂C have the same composition and crystal structure (fee),but M₁₂C has a smaller lattice constant than M₆C.The M₆C lattice constant is approximately 1.101 nm,whereas that of M₁₂C is approximately 1.085 nm.Chemical phase analysis and XRD detected an M₆C lattice constant ranging from 1.104to 1.106 nm in Alloy 2.Figure 7 shows bright-field image and the select area diffraction pattern (SADP) of particles in Alloy 2 after exposure at 750℃for 5000 h.Figure 7a shows a particle suspected to be M₆C phase with a size of500 nm,surrounded by sphericalγ'phase.SADP of the particle is marked with a red circle showing that it has fee structure and a lattice constant of 1.086 nm with a zone axis[110].Consequently,the particle could be identified as M₁₂C.

3.4 Evolution ofμphase

Theμphase is a topological close-packed (TCP) phases and shows a rhombohedral crystal structure.Butμphase and M₆C have a similar close-packed arrangement;hence,the alloy containing M₆C precipitate probably precipitatesμphase.When the total amount of molybdenum and tungsten in the alloy exceeds 10 wt%,it is reasonable to speculate the existence ofμphase [ 25, 26] .

Theμphase in Alloy 2 was analyzed by means of SEM,TEM,XRD and chemical phase analysis.Figure 8 shows SEM images of Alloy 2 after different exposure times at750℃.Figure 8a shows that after exposure for 1000 h,M₂₃C₆,M₆C and a small amount of needle-likeμphase precipitated around MC carbide particles.The amount of needle-likeμphase increased slightly after 3000-h exposure;however,a large amount ofμphase precipitated after5000 h.

Fig.7 a Bright-field image and b SADP of M₁₂C carbide after exposure at 750℃for 5000 h

Fig.8 SEM images of precipitates in Alloy 2 after exposure at 750℃for a 1000 h,b 3000 h and c 5000 h

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Table 3 Elemental component of different precipitates in Fig.9 with SEM+EDS analysis (wt%)

Table 3 lists elemental compositions of the granular M₆C precipitation and needle-likeμphase detected by EDS at different exposure times.Figure 9a shows that the granular precipitations (Position 1) have higher W content than the needle-like precipitation (Position 2),whereas the needle-like precipitations have lower Ni and Co content.The chemical phase analysis shows that the granular precipitation is M₆C phase,and the acicular precipitation isμphase.Therefore,M₆C has higher W content and lower Co and Ni content thanμphase.

A TEM micrograph of a needle-likeμphase is shown in Fig.10a.Theμphase precipitated in the grain is approximately 0.1μm in width and 6.5μm in length.Figure l0b is a local enlargement of the area shown in Fig.10a that contains many parallel fringes inμphase.Figure 10c shows the S ADP ofμphase.The diffraction spots stretched along the vertical direction of fringes indicate multiple stacking faults inμphase [ 27, 28] ,and the parallel fringes are the morphologies of stacking faults.Figure 10d shows a stacking fault crossing aγ'particle.Some needle-likeμphases precipitate randomly in matrix,as shown in Fig.10e.Someμphases nucleated around M₆C particles and grew outward.Some parallel fringes like stacking faults were observed nearμphase but identified as initialμphase morphology.

Because refractory W and Mo elements have larger atomic size and weight,their diffusion rates are very low;therefore,solution heat treatment did not completely eliminate the segregation of W and Mo.As a result,the content of refractory elements is much higher in some regions than in others.The elevated content leads to the formation ofγmatrix and causes these regions more easily reach super saturation state and induce M₆C and TCP phase precipitates [ 29] .Thermodynamic calculations show that the only phase that can precipitate from y matrix at 700and 750℃isμphase.Some reports have shown that M₆C is an unstable phase and a mesophase ofμphase.As M₆C has a similar atomic arrangement structure toμphase,it can transform intoμphase under specific thermodynamic conditions [ 30] .M₆C transforms into M₁₂C during 5000-h exposure at 750℃,but M₆C transforming intoμphase has not been observed during exposure,indicating that the formation ofμphase requires a higher thermodynamic driving force and precipitation temperature.

Because of the large particle size and the absence of a strengthening effect,the acicular precipitates deteriorate room-temperature plasticity and toughness [ 31] .Figure 11shows the evolved toughness of the two alloys exposed at700 and 750℃.During exposure at 700℃,the impact energy of both alloys decreases notably during the initial1000-h aging.The decrease trend slows between 1000 and5000 h and remains unchanged after 5000-h exposure.The impact absorbing energy of the two alloys decreases significantly during the initial 3000-h exposure and remains stable after 3000-h exposure at 750℃.

Fig.9 SEM images of precipitates in Alloy 2 after long-term exposure at 750℃:a,b 3000 h and b 5000 h

Fig.10 TEM images of stacking faults andμphase in Alloy 2 after exposure at 750℃for 5000 h:a,b needle-likeμphase;c S ADP ofμphase in a;d,e stacking faults andμphase

Fig.11 Impact absorbing energy after long-term exposure at a 700℃and b 750℃

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Table 4 Content of MC+M₆C+μprecipitations of Alloy 2 after exposure (wt%)

Because fracture behavior of the aged alloy is found to be intergranular,the distribution and quantity of carbides around the grain boundary have a decisive influence on the impact properties of the alloy.Because the MC phase distributed inside the grains remains stable during the longterm aging process,the effect of these particles on impact properties is negligible.The impact properties of the alloy do not decrease significantly with an increased amount of M₆C phases (Table 4 and Fig.11).Moreover,the dominant precipitation around the grains is confirmed to be M₂₃C₆.The larger amount and the shift in distribution from intermittent to continuous during the initial stage of aging reduce the impact absorbing energy significantly during the first 1000-h aging.

Alloy 2 has lower impact absorbing energy than Alloy 1because the higher W content in Alloy 2 increases the precipitation of M₂₃C₆,which in turn reduces the dislocation movement and results in more stress concentration during deformation.Consequently,the crack spreads more easily along the grain boundary.

The main precipitation period ofμphase at 750℃is between 3000 and 5000 h,but the impact absorbing energy was unchanged during this exposure time.Therefore,μphase has no distinct effect on the impact absorbing energy of Alloy 2.This may be becauseμphase mainly precipitates in grains,whereas the fracture of the alloy is mainly intergranular.

Table 4 shows the content of MC,M₆C andμphase in Alloy 2 after long-time exposure at 700 and 750℃.The values in Table 4 are the sum of the content of MC,M₆C andμphase,which have similar structures and cannot be separated by chemical phase analysis.During 700℃exposure,the content of M₆C and MC increases from 0.227to 0.353 after 1000 h and reaches 0.701 after 3000 h before remaining unchanged.MC is a stale phase and remains unchanged during exposure;therefore,the incremented precipitated phase content is M₆C.The content of the precipitated phases increases noticeably from 0.227 to2.058 after 3000 h at 750℃and reaches 2.894 after5000-h exposure.Figure 7 indicates that only MC,M₆C and a very small amount ofμphase precipitated after3000-h exposure,while the MC phase content does not change with increased exposure time.Thus,M₆C contributes the largest content of precipitates before 3000-h exposure,andμphase contributes more between 3000 and5000 h.

4 Conclusion

Thermodynamic calculation performed in the temperature range of 500-1500℃shows that the principal secondary phases in the alloy under equilibrium state consist of M₂₃C₆,γ'phase,MC,μphase and M₆C.The latter two phases were only found in the alloy with higher W content.The M₆C particles were found after solution heat treatment.During the subsequent 3000 h,the amount of this phase increased continuously with aging and then remained almost constant.The amount of precipitation at 750℃was much higher than that at 700℃.Part of M₆C transformed into M₁₂C after 5000-h exposure at 750℃.Needle-likeμphase precipitated after 1000-h exposure at 750℃,and the amount increased significantly during the following exposure time.The main elements of theμphase were Ni,W,Co,Cr and Mo.The impact energy of the alloy with high W content was low.However,μphase has no marked effect on the impact property during aging at 750℃.