J. XIAO et al.: EFFECT OF SOLUTION TREATMENT ON GAMMA-PRIME RECOVERY ... 671–678 EFFECT OF SOLUTION TREATMENT ON GAMMA-PRIME RECOVERY IN A CREEP-DAMAGED DIRECTIONALLY SOLIDIFIED Ni SUPERALLOY VPLIV RAZTOPNEGA @ARJENJA NA POPRAVO FAZE ’V DIREKTNO STRJENI IN ZARADI LEZENJA PO[KODOVANI Ni SUPERZLITINI Junfeng Xiao, Wenshu Tang * , Qing Nan, Sifeng Gao, Yongjun Li, Jiong Zhang Xi’an Thermal Power Research Institute Co., Ltd., 99 Yanxiang Road, Xi’an, China Prejem rokopisa – received: 2019-02-20; sprejem za objavo – accepted for publication: 2019-03-27 doi:10.17222/mit.2019.046 The effect of different solution conditions on the ’phase of a creep-damaged DZ411 superalloy during a process combining hot isostatic pressing (HIP) and re-heat rejuvenation was investigated. The analysis results showed that proper solution temperature and time are necessary for dissolving coarsened and rafted ’precipitates, allowing the optimum reprecipitation of the ’precip- itates during the following aging, avoiding incipient melting. It was also found that a higher solution temperature results in a smaller size and lower volume fraction of the ’precipitates. A two-hour solution treatment is enough to complete the dissolving of the remnant primary ’ precipitates and produce unimodal spheroidal secondary ’ precipitates. High cooling rates after the solution treatment give rise to finer and more secondary ’precipitates. The complete re-heat rejuvenation procedure including a full solution and two-step aging can produce duplex-sized ’precipitates including coarse cubic ’precipitates and fine spherical ’ precipitates, endowing the superalloy with excellent high-temperature creep properties. It was verified that the ’ precipitate morphology and creep properties of the creep-damaged DZ411 superalloy can be restored with a simple re-heat rejuvenation process. Keywords: directionally solidified superalloys, creep damage, re-heat rejuvenation, ’ precipitates Avtorji so raziskovali u~inek razli~nih raztopnih `arjenj na fazo ’ v vro~e izostatsko stiskani (HIP) superzlitini DZ411. Zlitina je bila po izdelavi med postopkom lezenja po{kodovana in renovirana (mikrostrukturno obnovljena) s postopkom ponovnega segrevanja. Rezultate so primerjali tudi z zlitino, ki ni bila HIP-ana. Rezultati analiz so pokazali, da je za raztapljanje grobih in ?naplavljenih? ’ izlo~kov potrebno raztopno `arjenje pri primerno visoki temperaturi in ustreznem ~asu zadr`evanja na tej temperaturi. To zagotavlja optimalno ponovno izlo~anje med naslednjim staranjem brez nataljevanja. Prav tako so ugotovili, da vi{ja temperatura raztopnega `arjenja povzro~a manj{e ’ izlo~ke in njihov manj{i volumski dele`. Dvourno raztopno `arjenje zadostuje za popolno raztapljanje primarnih ’izlo~kov in tvorbo enomodalnih krogli~nih sekundarnih ’izlo~kov. Ve~je hitrosti ohlajanja po raztopnem `arjenju povzro~ajo nastanek finej{ih in ve~je {tevilo sekundarnih ’ izlo~kov. Popolna prenova po{kodovane superzlitine s ponovnim segrevanjem vklju~uje popolno raztapljanje in dvostopenjsko staranje (izlo~evalno utrjevanje), kar povzro~a nastanek dvojne velikosti faze ’ z grobimi kubi~nimi in finimi krogli~nimi ’ izlo~ki. Tak{na mikrostrukturno prenovljena superzlitina ima odli~no visokotemperaturno odpornost proti lezenju. Avtorji so dokazali, da je mo`no morfologijo ’ izlo~kov in odpornost superzlitine DZ411 proti lezenju po njeni po{kodbi z lezenjem popolnoma restavrirati z enostavnim ponovnim procesom poprave oz. segrevanja. Klju~ne besede: direktno strjena superzlitina, po{kodbe zaradi lezenja, raztopno `arjenje. poprava, ’ izlo~ki 1 INTRODUCTION Turbine blades are one of the most important hot-sec- tion-path (HGP) components realizing energy transfor- mation in modern gas-turbine systems. Due to the com- plex service conditions including high temperature, high stress, corrosion and oxidation environment, turbine blades, after a long-term service exposure, inevitably suffer from microstructural damage, such as rafting or coarsening of ’ precipitates, all of which have a detri- mental effect on the blade performance. 1,2 Therefore, it is necessary to regularly replace or refurbish turbine blades. However, turbine blades require the application of superalloy materials, air cooling, thermal-barrier coat- ing and other advanced technologies. The complexity of the production process and the addition of high-cost re- fractory elements make the production cost of the blades higher and higher. The high cost of new replacement components provided a strong incentive for developing proper refurbishment procedures as economical methods to extend the service lives of blades. 3,4 Rejuvenation heat treatments, including a simple re-heat rejuvenation treatment or recovery cycles incor- porating both hot isostatic pressing (HIP) and heat treat- ment, were developed as a critical step in the refurbish- ment of degraded blades. 5,6 Previously, many studies about the recovery of the material microstructure in ther- mally exposed blades with an equiaxed structure had been conducted. 7–18 Subsequently, the HIP process was incorporated in the rejuvenation heat treatment of Inconel X-750 in 1977. 8 The early work consistently Materiali in tehnologije / Materials and technology 53 (2019) 5, 671–678 671 UDK 620.1:67.017:669.245:539.377 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 53(5)671(2019) *Corresponding author's e-mail: tangwenshu@tpri.com.cn showed that heat treatments alone could substantially re- covery the material microstructure and that an additional benefit was achieved through the HIP process for the re- juvenation cycle in some instances. 9,10 Although subse- quent studies reached mixed conclusions regarding the effectiveness of HIP recovery cycles, with some authors concluding that it is effective, 11,12 and others concluding that it is not, 13,14 there is still a common view that the HIP process has a positive impact on healing creep cavi- ties and casting shrinkage porosity, and that heat treat- ment can effectively improve the microstructure and properties. 15,16 At present, HIP recovery cycles are successfully used in the refurbishment and lifetime extension of degraded HGP components, applying a variety of equiaxed cast superalloys such as Udimet 500, Nimonic 115, Nimonic 263, IN738, GTD 111, Hastelloy X, FSX 414, and a few directionally solidified superalloys. 13,17–21 However, in the case of serviced components without any creep voids, HIP has no obvious benefit. 16 In fact, many directionally solidified superalloys are resistant to creep voids due to the minimum number of oriented grain boundaries, which indicates that they are likely to require the use of simple re-heat rejuvenation treatment consisting only of solution and aging treatments, which is a cheaper method for recovering the microstructure and properties of a directionally solidified superalloy. 16,22,23 For the re-heat rejuvenation treatment, the solution conditions intensely affect ’ precipitation characteristics. Mean- while, an improper solution treatment may cause unde- sirable effects, such as incipient melting or recrystalli- zation, especially for directionally solidified or single-crystal superalloys. To explore application possibility for simple re-heat rejuvenation treatment in recovering the microstructure of a service-damaged directionally solidified superalloy, we conducted a recovery study on the DZ411 superal- loy, 24 placing emphasis on the effects of different solu- tion conditions on the ’ precipitate characteristic of the creep-damaged DZ411 superalloy during a re-heat reju- venation process in this work. Finally, the recovery de- gree of the microstructure and creep properties of the creep-damaged DZ411 superalloy after a completed reju- venation procedure was evaluated. 2 EXPERIMENTAL PART Cylindrical bars of directionally solidified DZ411 superalloy, 12 mm in diameter and 200 mm in length, were produced under a vacuum atmosphere using the high-rate solidification method. 25 Before this, the DZ411 master alloy needed to be melted in a vacuum-induction furnace. The chemical composition of the superalloy was Ni-13.6Cr-9.14Co-4.9Ti-2.97Al-3.44W-1.6Mo-2.87Ta-0. 09C-0.01B (w/%), as determined with an Axios Max X-ray fluorescence spectrometer, iCAP6300 optical emission spectroscopy and AA600 atomic absorption. In the present investigation, after the heat treatment, the DZ411 superalloy was considered as the original super- alloy. Standard creep specimens of the DZ411 superalloy were taken from the original superalloy. Crystal orienta- tions of creep specimens were analyzed with the Laue X-ray back-reflection method, and crystal orientation de- viations were controlled within 10° from the [001] orien- tation. In order to simulate the creep damage of turbine blades during a long-term service, interrupted creep tests of the DZ411 superalloy were performed on an RDJ50 creep testing machine with a furnace attachment at 980 °C and 190 MPa. Interruption points at the secondary stage of creep were chosen based on creep-rupture curves. The DZ411 superalloy after the interrupted creep test was considered as the damaged superalloy in the present investigation. The damaged superalloy was then subjected to differ- ent heat-treatment cycles under an argon atmosphere in order to restore the degraded microstructure. The effects of the crucial solution-treatment parameters including temperature, holding time and cooling rate on the microstructure of the DZ411 superalloy were investi- gated. Detailed heat-treatment schedules are listed in Ta- ble 1. In order to prevent the superalloy samples from oxidizing during the heat-treatment process, the dam- aged superalloy was first sealed in a quartz tube filled with argon before it was place into the furnace. The metallographic samples were prepared by sec- tioning, grinding, polishing and etching with a solution of 10 g of copper sulfate in 50 mL of hydrochloric acid and 50 mL of water. The microstructures of all the sam- ples were examined with a PMG3 optical microscope (OM) and JSM-6460 scanning electron microscopy (SEM). Quantitative analyzing of the microstructures was conducted on the SEM micrographs using the Image Tool software as a semi-quantitative image analyzer. Since the quantitative metallographic investigation is principally a statistical measurement, many images of different conditions were analyzed. However, only one image from each set is shown in the present study as the representative one. Table 1: Re-heat treatment schedules for the creep-damaged DZ411 superalloy Cycle Full solution Partial solution Aging T (°C) t (h) Cool- ing T (°C) t (h) Cool- ing T (°C) t (h) Cool- ing A 1180 2 AC ------ B 1200 2 AC ------ C 1220 2 AC ------ D 1240 2 AC ------ E 1220 1 AC ------ F 1220 3 AC ------ G 1220 4 AC ------ H 1220 2 FC ------ I 1220 2 AC 1121 2 AC 843 24 AC Note: AC – air cooling, FC – furnace cooling J. XIAO et al.: EFFECT OF SOLUTION TREATMENT ON GAMMA-PRIME RECOVERY ... 672 Materiali in tehnologije / Materials and technology 53 (2019) 5, 671–678 3 RESULTS AND DISCUSSION 3.1 Original and creep-damaged microstructures The microstructures of the original and damaged superalloy are shown in Figure 1. As can be observed in Figure 1a, the microstructure of the original superalloy exhibits a bimodal distribution of coarse primary ’ pre- cipitates and fine secondary ’ precipitates within the austenitic matrix. The primary ’ precipitates, with the average size of 0.6 μm and a volume fraction of 50.3 %, have a cubic morphology, while the secondary ’precipi- tates, with the average size of 0.11 μm and a volume fraction of 2.49 %, have a spherical morphology. How- ever, the ’ precipitates of the damaged superalloy lost their original morphology (as shown in Figure 1b).I n particular, the primary ’ precipitates are coarsened and converted into spheroidal ones at the expense of the fine secondary ’ precipitates, resulting in a sharp decrease in the secondary ’precipitates. It is believed that the coars- ening of the ’ precipitates in nickel-based superalloys is in accordance with the Ostwald ripening mechanism, ac- cording to the Lifshitz-Slyozov-Wagner theory and sub- sequent modifications, 26 and that the coarsening process of the primary ’ precipitates is driven by the reduction in the total interfacial energy. Table 2 summarizes the statistical data of the ’ precipitate characteristics for dif- ferent states of the superalloy. It can also be seen that the volume fraction of the primary ’ precipitates of the damaged superalloy decreased significantly compared with that of the original superalloy. J. XIAO et al.: EFFECT OF SOLUTION TREATMENT ON GAMMA-PRIME RECOVERY ... Materiali in tehnologije / Materials and technology 53 (2019) 5, 671–678 673 Figure 2: Optical micrographs of the damaged superalloy after solution treatments at: a) 1180 °C, b) 1200 °C, c) 1220 °C and d) 1240 °C Figure 1: ’ precipitate microstructure for the original and damaged superalloy 3.2 Microstructure after various solution-treatment schedules To determine the proper solid solution temperature, various solution treatments were initially performed on the damaged superalloy at four holding temperatures (1180, 1200, 1220 and 1240) °C for2ha n dsubse- quently cooled in air to room temperature (correspond- ing to heat-treatment cycles A–D), respectively. Optical micrographs of the superalloy after the solution treat- ments at different holding temperatures are shown in Figure 2. It is clearly seen that the size of - ’ eutectic decreased as the solution temperature increased from 1180 °C to 1240 °C. However, when the solution treat- ment was performed at 1240 °C, incipient melting in the interdendritic region occurred and the - ’ eutectic phase almost disappeared. Therefore, 1240 °C is the solidus temperature of the damaged superalloy. Figure 3 shows the ’ precipitate microstructure of the damaged superalloy after the solution treatment at different holding temperatures (1180, 1200, 1220 and 1240) °C for 2 h and subsequent air cooling to room temperature (heat treatment cycles A–D, respectively). The measured size and volume fraction of the ’ precipi- tates in the damaged superalloy before and after different heat-treatment cycles are presented in Table 2.Itisobvi- ous that the obtained ’ precipitates, except for heat- treatment cycle A, lost their rafting and coarsening char- acteristics. As seen in Figure 3a, it is clear that the 1180 °C/2 h/AC solution treatment (heat-treatment cycle A) gave rise to a decrease in the size and volume fraction of the ’ precipitates and produced a nearly bimodal, spheroidal ’ microstructure with coarse and fine ’ pre- cipitates within the austenitic matrix, compared with the ’ precipitates in the damaged superalloy. In fact, 1180 °C is the subsolvus temperature for the ’ precipi- tate dissolution in the DZ411 superalloy. It indicates that a partial dissolution of the coarsened prime ’ precipi- tates in the damaged superalloy occurred. A large num- ber of undissolved prime ’precipitates, namely the rem- nant prime ’ precipitates, are retained in the matrix. Meanwhile, a partial dissolution of the prime ’ precipi- tates incorporated more solute atoms into the matrix, leading to the formation of fine secondary ’precipitates, namely cooling precipitates, in the space between the remnant ones. Hence, it can be concluded that the coarse prime ’ precipitates are the remnant prime ’ precipi- tates and the finer ones are the secondary ’ precipitates that formed during the cooling process after the solution treatment. However, when the solution temperature increased up to 1200 °C (heat-treatment cycle B), the ’ precipitates lost the entire rafted and coarsened morphology and ex- hibited a smaller size and a nearly unimodal, spheroidal microstructure (shown in Figure 3b), compared with the ’ precipitates in the damaged superalloy. The formation of unimodal, spheroidal ’ precipitates implies that al- most a complete dissolution of the remnant prime ’ pre- cipitates occurred after heat-treatment cycle B. It can be deduced that the formed unimodal, spheroidal ’ precipi- J. XIAO et al.: EFFECT OF SOLUTION TREATMENT ON GAMMA-PRIME RECOVERY ... 674 Materiali in tehnologije / Materials and technology 53 (2019) 5, 671–678 Figure 3: ’precipitate microstructure of the damaged superalloy after different solution temperatures and cooling rates: a) Sample A (1180 °C/2 h, air cooling), b) Sample B (1200 °C/2 h, air cooling), c) Sample C (1220 °C/2 h, air cooling), d) Sample D (1240 °C/2 h, air cooling) tates are almost the secondary ’ precipitates. As the so- lution temperature further increased up to 1220 °C (heat-treatment cycle C) and 1240 °C (solution-heat treatment cycle D), the secondary ’ precipitates were still nearly unimodal and spheroidal, exhibiting a slight decrease in the size and an increase in the volume frac- tion (see Figures 3c and 3d and Table 2). It is believed that further dissolution of the remnant prime ’ precipi- tates at a higher solution temperature incorporated more solute into the matrix, resulting in finer secondary pre- cipitates being formed during the cooling process after the solution treatment. In view of the fact that the solu- tion treatment at 1240 °C caused incipient melting, the suitable solution temperature is determined to be 1220 °C for the re-heat rejuvenation treatment of the creep-damaged DZ411 superalloy. Besides the solid-solution temperature, the holding time and cooling rate are also important factors, influ- encing the ’ precipitate characteristics. The ’ precipi- tate microstructure of the damaged superalloy after the solution treatment at 1220 °C for different holding times (1, 3 and 4) h and the subsequent air cooling to room temperature (corresponding to heat-treatment cycles E–G) are shown in Figure 4a to 4c, respectively. Com- pared with the ’ precipitates obtained after heat-treat- ment cycle C, it is seen that as the holding time increased from1hto2h,allthe ’ precipitates exhibited a similar unimodal, spheroidal microstructure and also a slight in- crease in the volume fraction, indicating that the holding time of1hatthesolution temperature of 1220 °C is not enough to dissolve the remnant prime ’ precipitates. However, as the holding time further increased up to 3 h (heat-treatment cycle F) and 4 h (heat-treatment cycle G), the ’ precipitates showed a slight increase in the av- erage size and squareness, and also a slight decrease in the volume fraction, possibly resulting from the contri- bution of the elastic strain energy during the solution treatments. 27 Figure 4d represents the ’ precipitate microstructure produced after heat-treatment cycle H, in which the cooling from the solution temperature to room temperature is performed by the furnace. Comparing the ’ precipitate microstructures ob- tained after heat-treatment cycles C and H, it can be eas- ily found that heat-treatment cycle H produces a lower density and large cubes of the ’ precipitates, unlike the finer, unimodal, spheroidal ’ precipitates obtained after heat-treatment cycle C (Figure 3c). This can be related to the cooling rate during the cooling process after the solution treatment. During the furnace cooling, the nu- cleation rate of the ’ precipitates is expected to be lower; meanwhile, a longer time is available for the dif- fusion of solute atoms and growth of the precipitates due to the low cooling rate. 28 However, according to the sta- tistical parameters for the ’ precipitates after cycles C andH( Table 2), it is to be noted that the volume frac- tions of the ’ precipitates after cycles C and H are al- most the same, indicating that the ’ precipitates grow slowly due to the solute ejection from the matrix during slow cooling, partially compensating for the advantage of the high ’ nucleation rate caused by air cooling. J. XIAO et al.: EFFECT OF SOLUTION TREATMENT ON GAMMA-PRIME RECOVERY ... Materiali in tehnologije / Materials and technology 53 (2019) 5, 671–678 675 Figure 4: ’ precipitate microstructure of the damaged superalloy after different solution times and cooling rates: a) Sample E (1220 °C/1 h, air cooling), b) Sample F (1220 °C/3 h, air cooling), c) Sample G (1220 °C/4 h, air cooling), d) Sample H (1220 °C/2 h, furnace cooling) Table 2: Statistical data for the ’precipitate characteristic parameters for the damaged and rejuvenated superalloy Condition ’ type Average size /μm Volume fraction /% Original superalloy Primary ’ Secondary ’ 0.600±0.029 0.105±0.022 50.325±0.221 2.486±0.006 Damaged superalloy Primary ’ Secondary ’ 0.800±0.032 - 37.495±0.462 - A Primary ’ Secondary ’ 0.696±0.079 0.101±0.025 25.221±0.816 11.442±0.007 B Primary ’ Secondary ’ - 0.292±0.032 - 39.480±0.020 C Primary ’ Secondary ’ - 0.154±0.013 - 40.140±0.023 D Primary ’ Secondary ’ - 0.147±0.026 - 41.95±0.013 E Primary ’ Secondary ’ - 0.151±0.019 - 38.587±0.018 F Primary ’ Secondary ’ - 0.162±0.028 - 40.931±0.282 G Primary ’ Secondary ’ - 0.169±0.017 - 37.854±0.281 H Primary ’ Secondary ’ - 0.586±0.049 - 40.966±0.416 I Secondary ’ Tertiary ’ 0.394±0.042 0.071±0.003 42.510±0.200 6.488±0.004 3.3 Microstructure and creep properties of the rejuve- nated superalloy Due to a lack of a complete rejuvenation procedure, a single solution treatment could not fully recover the microstructure of the damaged superalloy, achieving its original state. Figure 5 shows the ’ precipitate micro- structure of the damaged superalloy after a complete re-heat rejuvenation procedure consisting of a full solu- tion and two-step aging at the following conditions: 1220 °C/2 h/AC + 1121 °C/2h/AC + 843 °C/24 h/AC (heat treatment cycle I). It is obvious that heat-treatment cycle I produced a duplex-size ’ precipitate micro- structure with coarse cubic ’precipitates and fine spher- ical ’ precipitates, which are analogous to those of the original superalloy except for the smaller ’ precipitates. Especially the size of the coarse ’ precipitates is larger and the size of fine ’ precipitates is smaller than that of the ’ precipitates obtained after heat-treatment cycle C. In fact, the aging temperature of 1121 °C is below the temperature of the ’ precipitate dissolution in the DZ411 superalloy. So, it is concluded that the obtained coarse ’ precipitates were derived due to the growth of the secondary ’ precipitates, which formed during the cooling after the solution treatment and the fine ones are tertiary ’ precipitates which reprecipitated during the aging treatment. This kind of ’ precipitate microstructure is expected and theoretically desired as the most optimized microstructure that can provide for very good high-tem- perature properties. Figure 6 shows creep curves for the original, damaged and rejuvenated superalloy at 980 °C and 190 MPa. It is evident that the creep curve of the DZ411 superalloy after heat-treatment cycle I is remark- ably similar to that of the original superalloy. Moreover, the creep-rupture time of the rejuvenated superalloy is near to that of the original superalloy, but about four times longer than that of the damaged superalloy. There- fore, it is feasible to restore the microstructure and creep properties of directionally solidified DZ411 superalloy with a simple re-heat rejuvenation treatment. 4 CONCLUSIONS A creep-damaged DZ411 superalloy was initially ob- tained with an interrupted creep test. The damaged superalloy was then subjected to different solution-treat- ment cycles and a two-step aging treatment to reveal the influences of different solution conditions on ’ precipi- tate characteristics and investigate the evolution of de- generated ’ characteristics. Based on the observations and analysis results, the following conclusions were ob- tained: 1. The creep test gave rise to a microstructural degra- dation of the original superalloy. The microstructural de- J. XIAO et al.: EFFECT OF SOLUTION TREATMENT ON GAMMA-PRIME RECOVERY ... 676 Materiali in tehnologije / Materials and technology 53 (2019) 5, 671–678 Figure 6: Creep curves for the original, damaged and rejuvenated superalloy at 980 °C and 190 MPa Figure 5: ’ precipitate microstructure of the rejuvenated superalloy terioration was evidenced mainly by the coarsening and rafting of prime ’ precipitates and disappearance of sec- ondary ’ precipitates. 2. Proper solution conditions including 1220 °C and 2 h provide for a dissolution of the coarsened and rafted prime ’ precipitates and reprecipitation of fine second- ary ’ precipitates during the cooling process after the solution treatment, while avoiding incipient melting dur- ing the re-heat rejuvenation treatment of the damaged superalloy. 3. Solution treatment at a higher temperature results in a smaller size and higher volume fraction of the ’pre- cipitates. However, solution treatment at a longer holding time produces a larger size and a lower volume fraction of the ’ precipitates. High cooling rates after the solu- tion treatment give rise to finer and more secondary ’ precipitates. 4. A complete rejuvenation procedure including a full solution and two-step aging can produce a duplex-sized ’precipitate microstructure with coarse cubic ’precipi- tates and fine spherical ’ precipitates, providing very good high-temperature creep properties. 5. It is verified that the microstructure and creep properties of the creep-damaged DZ411 superalloy can be restored with a simple re-heat rejuvenation treatment. Acknowledgment This paper was supported by the National Natural Science Foundation of China (No. 51601145). 5 REFERENCES 1 J. Wang, L. Z. Zhou, L. Y. Sheng, J. T. 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