Effects of Precipitates and Oxide Dispersion on the High-temperature Mechanical Properties of ODS Ni-Based Superalloys

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ISSN 1225-7591(Print) / ISSN 2287-8173(Online)

Effects of Precipitates and Oxide Dispersion on the High-temperature Mechanical Properties of ODS Ni-Based Superalloys

GooWon Noh

^a,b

, Young Do Kim

^b

, Kee-Ahn Lee

^c

and Hwi-Jun Kim

^a,

*

a

Liquid processing & Casting R&D Group, Korea Institute of Industrial Technology, Incheon 406-840, Republic of Korea

b

Department of Materials Science and Engineering, Hanyang University, Seoul 04763, Republic of Korea

c

Department of Materials Science and Engineering, Inha University, Incheon 22212, Republic of Korea (Received February 17, 2020; Revised February 24, 2020; Accepted February 24, 2020)

...

Abstract In this study, we investigated the effects of precipitates and oxide dispersoids on the high-temperature mechanical properties of oxide dispersion-strengthened (ODS) Ni-based super alloys. Two ODS Ni-based super alloy rods with different chemical compositions were fabricated by high-energy milling and hot extrusion process at 1150

^o

C to investigate the effects of precipitates on high-temperature mechanical properties. Further, the MA6000N alloy is an improvement over the commercial MA6000 alloy, and the KS6000 alloy has the same chemical composition as the MA6000 alloy. The phase and microstructure of Ni-based super alloys were investigated by X-ray diffraction and scanning electron microscopy. It was found that MC carbide precipitates and oxide dispersoids in the ODS Ni-based super alloys developed in this study may effectively improve high-temperature hardness and creep resistance.

Keywords: Ni-base-ODS, Precipitates, Creep property, Hf-oxide, Mechanical alloying

...

1. Introduction

Oxide dispersion strengthened (ODS) Ni-based super alloys have higher available temperature and excellent high temperature properties compared to conventional super alloys [1, 2]. Therefore, various studies have been conducted on ODS Ni-based super alloys with excellent high-temperature characteristics for industrial fields such as aerospace, defense, and gas turbine applications. The oxide dispersion strengthening mechanism of various strengthening methods plays the most important role to provide superior strength and stability at elevated tem- peratures, since dispersed oxide particles interrupt move- ment of dislocation effectively at high temperature via the “Pinning effect” phenomenon. It is well known that the “Pinning effect” depends on the density and size of precipitates and oxide dispersoids [3].

ODS Ni-based super alloys are commercially available as MA6000 and MA758 alloys. The chemical composi-

tion of MA6000 and MA758 alloys are Ni-15Cr-4.5Al- 2.5Ti-2Mo-4W-0.15Zr-1.1Y

₂

O

₃

and Ni-30Cr-0.3Al-0.5Ti- 1.0Fe-0.05C-0.6Y

₂

O

₃

, respectively. Considering these gen- eral oxide dispersion strengthening mechanisms, the strengthening mechanism of Ni-based ODS alloys can be represented by the following three strengthening mecha- nisms: 1) solid solution hardening (γ-phase) by the high solubility of Cr in Ni (47 wt.% at 1618 K), 2) precipita- tion hardening by Ni

₃

(Al,Ti) precipitate (γ’-phase), and 3) dispersion hardening through the formation of oxide dis- persoid-based Y-metal-O phase. An alloying method with a good combination of precipitation hardening and dis- persion hardening was developed for the first time at the International Central Laboratory (INCO) in the United States by a mechanical alloying (MA) process [4]. The purpose of the MA process is to simultaneously alloy an oxide with a high melting point and a metal having a rel- atively lower melting point, which cannot be alloyed by the general melting process, and then to finely and uni-

- 노구원: 학생, 김영도·이기안: 교수, 김휘준: 수석연구원

*Corresponding Author: Hwi-Jun Kim, TEL: +82-32-850-0406, FAX: +82-32-850-0390, E-mail: [email protected]

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tive elements, which have no solubility in Ni, and fabri- cated alloy powders by a gas atomization process.

Through this study, we attempted to obtain a stable car- bide and oxide at high temperature to improve the mechanical properties and minimize the disadvantages of ODS alloy manufacturing method using a conventional MA process. In addition, we propose a progressive liq- uid sintering (PLS) process, which is a new heat treat- ment method. The three-step heat treatments after hot extrusion must be performed in order to obtain the γ’- phase, which is the main hardening phase of Ni-based ODS alloys. However, this method also has a disadvan- tage that the manufacturing time is considerably long; the manufacturing time is reduced through the PLS process.

The described method was effective at removing pores inside alloy material and suppress the diffusion of solute through the liquid phases having a partial melting state with a fraction of 20 % to 80 %. Therefore, we report on a new ODS alloy fabrication process, and discuss the effects of stable precipitates and dispersoids on the result- ing high temperature properties at high temperatures.

2. Experimental procedure

ODS alloy powders were prepared by high-vacuum gas atomization (HVGA) in a vacuum atmosphere in order to investigate the variation of high temperature characteris- tics by chemical composition. The nominal chemical

powders were mechanically alloyed for a total of nine hours in an oxygen atmosphere by repeating the two-step condition (first step: 480 RPM for 50 minutes, second step: 245 RPM for 10 minutes). Phase analysis was per- formed using XRD (model: X’ Pert-Pro MPD/PANalyti- cal) to observe phase changes before and after MA. A steel can was prepared for hot extrusion of the milled powders. The milled powder was poured into the steel can and was degassed at 773 K for six hours under a high vacuum atmosphere. Subsequently, the can was extruded at an extrusion ratio of 16:1 at 1423 K. The extruded can is rod shaped with Φ 16. These rods were post-treated under an argon atmosphere at a moving speed of 70 mm/min at about 1473 K by the PLS method.

The high temperature mechanical properties like hard- ness and tensile strength were measured up to 1253 K.

The creep properties were evaluated at a stress of 135 MPa and creep temperature of 1253 K. After the creep tests, the microstructure was observed by field emission scanning electron microscopy (FESEM) and energy dis- persive spectrometer (EDS) (TESCAN MIRA3).

3. Results and Discussion

Figure 1 is the EPMA image of the MA6000N alloy manufactured by gas atomization. Figure 2 shows a typi- cal secondary electron (SE) image with dendrite obtained from gas atomized metal powders. The Y, Hf, and W ele-

Table 1. Nominal compositions of Ni super alloys [Wt.%]

Ni Cr Al Ti Mo W Ta Zr B Nb Y C Hf Y

₂

O

₃

MA6000N Bal. 15 2.5 3 2 4 2 0.2 0.1 2 0.8 0.1 2 0.5

KS6000 Bal. 15 4.5 2.5 2 4 2 0.15 0.01 - - - - 1.1

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ments in MA6000N powder were homogeneously distrib- uted at the interdendritic regions by EDS mapping. On the powder surface, Y and O elements were observed to coexist, because Y has an exceptionally high affinity for oxygen.

In the conventional MA process, a nano-sized Y

₂

O

₃

with MA6000 alloy powder was added at 1.1 wt. %.

However, the MA600N was milled by addition of a min- ute amount of 0.5 wt. %, because there is a Y

₂

O

₃

layer on the surface. We utilized XRD to identify phase varia- tion due to the addition of Y

₂

O

₃

and MA results. Figure 2 shows the XRD results of MA6000N powder before (a) and after (b) ball milling (MA). The XRD result before ball milling indicated that the MA6000N alloy powder is composed of α-Ni, Y

₂

O

₃

, and Ni

₅

Y phases. Therefore, when the EPMA and XRD results are compared, it can be seen that the surface oxide is Y

₂

O

₃

and the internal

phases are α-Ni and Ni

₅

Y. The XRD result after ball mill- ing showed a broad pattern and MC carbide phase. The MC carbide within a Ni-based super alloy tends to be degraded by M

₂₃

C

₆

or M

₆

C carbide during ball milling of heat treatment.

Hardness variations of hot-extruded rods were investi- gated at various temperatures from room temperature to 1273 K after hot extrusion. Figure 3 displays the result of hardness variation for MA6000N and KS6000. As the measurement temperature increased, the hardness of both alloys decreased gradually. After 1073 K, the hardness dropped sharply from the previous zones [7-9]. In addi- tion, the hardness of the MA6000N rod was higher than the KS6000 rod throughout all zones (zone 1, zone 2, and zone. 3). Dispersion strengthening affects mechani- cal properties at high temperatures due to particle size and type, spacing, and dispersion. It is also known that Fig. 1. A cross-sectional SEM image and the corresponding EPMA elemental maps of MA6000N atomized powder.

Fig. 2. XRD patterns of (a) MA6000N atomized powder and (b) MAed powder.

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dispersoids, which were finely and homogeneously dis- persed, play a role in improving mechanical properties at both room temperature and high temperature.

Microstructure was observed to investigate the effect of these strengthening factors on MA6000N alloys. Figure 4

ness at 1273 K can be estimated for the following rea- sons [9-12]. The main strengthening phase (γ' phase with a face-centered cubic (FCC) crystal structure and coher- ent relationship with matrix) is greatly influenced by the mechanical properties at high temperature [13]. How- ever, the hardness value is drastically reduced at tempera- tures above 1073 K, because the γ' phase is dissolved into the matrix.

Fig. 3. Vickers' hardness values at various temperatures.

Fig. 4. FESEM and EDS of MA6000N after PLS at 1473 K.

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Figure 5 shows the creep test results of the MA6000N rod heat-treated by the PLS process. The sample finally failed after 650 hours at a constant load and temperature of 135 MPa and 1253 K, respectively. It is superior to the results reported in literatures tested at constant load and temperature of 230 MPa and 1253 K, respectively [14].

In addition, the Larson-miller index P

₁

of MA6000N at 1273 K was 3.3 × 10

⁴

, and Mino, K. showed that P

₁

of Alloy98 alloy was 3 × 10

⁴

under the same experimental conditions. Samples of MA6000N are expected to break after 264.85 hours with 135 MPa stress at 1273 K and after 1.148 hours with Alloy98 alloy [15]. The improve- ment of creep resistance of the ODS alloys in the present study can be contributed by solid solution strengthening, precipitation hardening, dispersion strengthening, direc- tionally grain growth by secondary recrystallization by preventing intergranular slip, and minimizing transverse rupture. Figure 6 shows the EBSD results of a MA6000 rod heat-treated by the PLS process. It was confirmed that the grains did not directionally grow and had a ran- dom orientation relationship.

Therefore, it can be seen that our sample depends on

the mechanical properties only by γ' precipitate and car- bide-based dispersoid. However, the volume fraction of γ'-phase in commercial alloys, INCO MA6000, is about 55%, while the MA6000N alloy is considerably less at about 7.66% (Table 2).

Thus, the hardening result of MA6000N alloys is expected to be due to the dispersion strengthening mech- anism via the insoluble second phase, such as Hf-oxides, nano-sized M

₂₃

C

₆

carbide, rather than the γ'-phase in Zone 3 (Table 3).

The cavitation fraction tends to increase in the region close to the fractured surface. However, the fracture fea- tures were intergranular with no specific direction, unlike the trans-granular fracture where the crack propagates in the 45° direction. In other words, the MA6000N alloy has a creep similar to that of INCO MA6000 alloy, which Fig. 5. Creep behavior of MA6000N at 1253 K. Fig. 6. EBSD of MA6000N after PLS at 1473 K.

Table 2. The statistical area fractions and the average sizes of dispersoids and precipitates in the grain interior and along the grain boundaries

MA6000N alloy γ’ Carbide

Area Fraction (%) 7.66 5.06

Table 3. Temperature effects on strengthening mechanisms Effective

Temperature Strengthening

mechanism

Zone. 1

~ 0.3 T

_m

Zone. 2

~ 0.6 T

_m

Zone. 3

~ 0.9 T

_m

Comments

Work Hardening O X X Dislocation/Dislocation Interaction

Grain Size Refinement O X Inverse proportion Grain Boundary Area-Dislocation Interaction

Solid Solution Strengthening O X Lattice Strain Fields: Modulus Change, Dislocation interaction

Precipitation Strengthening O O Metastable 2

^nd

Phase Particles, impede Dislocation motion

Dispersion Strengthening O O O Insoluble 2

^nd

Phase, stabilize grain, subgrains and creep

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is a commercially available alloy, because it prevents grain boundary sliding due to Hf-based oxide disper- soids and M

₂₃

C

₆

carbide, which are finely present in grain boundaries despite the intergranular fracture. The mechanism for creep characteristics needs further analy- sis and will be reported in subsequent investigations.

4. Conclusion

1) In this study, we proposed a new alloy system and process with similar performance to the conven- tional ODS alloys while reducing the milling time and increasing productivity.

2) The main hardening mechanisms were dispersion hardening by γ'-phase and carbide-based M

₂₃

C

₆

car- bide. Both phases were high in brittleness, so the existing 48-hour milling time could be shortened to nine hours.

3) The tertiary creep before creep rupture was rapidly deformed, and then fractured because the fraction of work hardening is larger than the recovery, unlike the secondary creep, and the vacancy increases rap- idly at higher temperatures. The effect of Hf-oxides and M

₂₃

C

₆

carbide on mechanical properties is greater than γ'-phase in our alloy system. Therefore, it is expected that the volume fraction of the γ'-phase dissolved in the matrix is less than other commer- cial alloy systems and the creep characteristics are excellent.

References

8[1] W. Betteridge and J. Heslop: Nimonic Alloys and Other Nickel-base High Temperature Alloys 2nd ed., J. Heslop (Ed.), Originating Research Org., United Kingdom, (1974).