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Fundamental Understanding of Creep-Fatigue Interactions in 9Cr- 1Mov Steel Welds

Fundamental Understanding of Creep-Fatigue Interactions in 9Cr- 1Mov Steel Welds

Project No. 13-4948

Fundamental Understanding of -Fatigue Interactions in 9Cr- 1MoV Welds

Reactor Concepts Research Development and Demonstration (RCRD&D)

Wei Zhang The Ohio State University

William Corwin, Federal POC Meimei Li, Technical POC Final Project Report Revision 1

DOE Award Number: DE-NE0000708

Recipient name: The Ohio State University

Project title: Fundamental Understanding of Creep-Fatigue Interactions in 9Cr-1MoV Steel Welds

Principal investigators: PI: Co-PI: Wei Zhang Michael J. Mills Professor Professor [email protected] [email protected] 614-292-0522 614-643-3463

Graduate students Dr. H. Collin Whitt (Ph.D. in and Engineering, 2019) Mr. Tyler Payton (MS in Welding Engineering, 2017) Mr. Kaiwen Zhang (Ph.D. Candidate) The Ohio State University

Research collaborators Drs. Yanli Wang and Zhili Feng, Oak Ridge National Laboratory

Project period start and end From 1/17/2014 to 7/15/2018 date: Report submission date: 5/30/2019

Final Project Report for DE-NE0000708 - Revision 1

Fundamental Understanding of Creep-Fatigue Interactions in 9Cr-1MoV Steel Welds

Submitted by H. Collin Whitt, Kaiwen Zhang, Michael J. Mills and Wei Zhang The Ohio State University

Table of Contents a. Abstract and Executive Summary ...... 4 b. Introduction and Objectives ...... 7 c. Background ...... 9 c.1 Creep-strength Enhanced Ferritic ...... 9 c.2 Alloy Development for Elevated Environments ...... 11 c.3 Creep-fatigue: A More Accurate Description of in G91 Components ...... 14 c.4 Weldability Issues Associated with 9-12% Cr CSEFS ...... 18 c.5 The Future of CSEFS ...... 20 c.6 9-12% Cr CSEFS: Summary and Unanswered Questions ...... 22 d. Experimental Procedures ...... 23 d.1 Materials and Welding ...... 23 d.5 Metallography ...... 26 d.6 Characterization ...... 26 e. G91 Base Characterization and Creep-Fatigue Studies ...... 29 e.1 Introduction and Objectives for Base Metal CF Studies ...... 29 e.2 Experimental Procedures ...... 31 e.3 Results and Discussion ...... 32 e.4 Summary and Conclusions of Base Metal CF Studies ...... 40 f. Creep-fatigue Deformation and Characterization of G91 Weldments ...... 41 f.1 Introduction and Objectives for G91 Weld Studies ...... 41 f.2 Materials and Experimental Procedures ...... 43 f.3 Results and Discussion ...... 44 f.4 Local Strain Accumulation in HAZ ...... 50

2 f.5 Summary and Conclusions for G91 Weld Studies ...... 51 g. Overall Summary, Conclusions and Future Work ...... 52 g.1 G91 Base Metal Creep-fatigue Studies ...... 52 g.2 G91 Low heat-input Welding and Creep-fatigue Studies ...... 53 g.3 Overall Conclusions ...... 54 h. List of publications ...... 54 i. Acknowledgements ...... 56 j. References ...... 56 Figure Captions ...... 63

3 a. Abstract and Executive Summary Ferritic-martensitic steels constitute a key class of materials for power generation due to their creep-resistant properties. The current study examines the creep and creep-fatigue (CF) properties of a workhorse ferritic-martensitic steel alloy, Grade 91 or G91 (9Cr-1MoV). This steel is commonly used in boiler and piping applications at elevated . While base metal properties are important, component failure most commonly occurs in the welded region of components in-service. Therefore, the present work consists of studies on base metal and welded ferritic-martensitic components, examining the effect of CF deformation and weldment processing on microstructure and elevated temperature performance. First, creep-fatigue properties of G91 base metal are examined. In service, G91 components are subjected to elevated temperatures and cyclic stresses, leading to the accumulation of creep-fatigue damage. There is a limited data on creep-dominated CF damage of G91 base metal in the current literature. This study examines the creep-fatigue behavior of G91, including mechanical response under various -controlled testing conditions at 600C and 650C. The testing is done using a custom-built, load-controlled creep-fatigue testing apparatus at OSU. Microstructural studies utilize techniques including: scanning electron microscopy (SEM), scanning transmission electron microscopy (STEM), electron backscattered diffraction (EBSD) and transmission Kikuchi diffraction (TKD). Quantitative microstructural studies track substructure coarsening and density as a function of creep-fatigue deformation. Significant anelastic backflow is observed at minimum load during every creep- fatigue test conducted. The effects of loading parameters on creep-fatigue rupture life and anelastic backflow are also studied. The differences between monotonic creep and creep- fatigue, which lead to accelerated failure under creep-dominated CF deformation are examined. Second, creep-fatigue properties of G91 weldments are also assessed by studying a conventional flux-cored arc welding process (FCAW) as well as a non-conventional cold metal transfer (CMT) welding process. Cross-weld specimens from each weldment were creep-fatigue deformed under multiple loading conditions at 650C. Ruptured weldments are examined using characterization techniques including SEM and STEM diffraction-contrast imaging (STEM-DCI). Specimens welded using the CMT process significantly outperform the FCAW weldments. Further characterization reveals that changes in precipitate size and distribution as well as differences in subgrain size and dislocation density between the two welding processes result in the differences in creep-fatigue strength.

The key apparatus developed or utilized in the project included:  Developed a cost-effective dwell CF testing system based on a standard lever-arm creep-testing frame with a load train featuring a pneumatic actuator. The system was equipped with Linear Variable Differential Transformer (LVDT) extensometers as well as a digital acquisition system to record strain, temperature, dwell time and loading cycles.

4  Developed a high-temperature digital image correlation (DIC) based measurement system to map the strain localization in the Grade 91 weld in-situ during CF testing.  Developed the welding procedure for joining Grade 91 steel plates using a low heat input arc welding process based on cold metal transfer (CMT).  Utilized Gleeble®, a thermal-mechanical testing system, to produce specimens with simulated ICHAZ microstructure, which were subsequently used in CF testing to isolate the effect of HAZ microstructure on CF properties.  Utilized high-resolution electron microscopy to characterize subgrain size, dislocation density, and precipitate distribution of Grade 91 steel and welds.

The overall tasks and results obtained in the project included:  Fabricated welded joints of G91 plates using the conventional FCAW and low heat input CMT processes.  Generated a large amount of CF data of G91 base metal and cross-weld specimens tested under various load-controlled (thus creep-dominated) testing conditions.  Performed extensive microstructural characterization including evolution of subgrain size, dislocation density, and precipitate size distribution during CF testing of Grade 91 base metal.  Established correlation between observed anelastic backflow of strain at minimum load and CF deformation.  Measured evolution of local HAZ strain in cross-weld specimens in-situ using high- temperature DIC.  Performed characterization of location-specific ICHAZ microstructure for subgrain size, dislocation density, and precipitate size distribution before and after CF testing.  Performed computational simulation of carbon diffusion during welding and subsequent CF testing.

By studying the mechanical response to CF loading coupled with advanced microstructure characterization, the following key conclusions were drawn:

Base metal of Grade 91 steel:  After , G91 base metal contains a high dislocation density, refined subgrain size and fine distribution of particles along low angle grain boundaries (LAGBs) and low angle grain boundaries (HAGBs), all of which continuously evolve during CF deformation.

5  Studying microstructural evolution reveals that subgrain coarsening is sluggish at the onset of CF damage but occurs rapidly in later stages of CF deformation. Cooperatively, dislocation density decreases rapidly at the onset of CF damage and decreases at a continuously slower rate as a function of CF cycling.  A significant amount of anelastic backflow of strain occurs during every load-controlled CF cycle. Anelastic backflow increases with time spent at minimum and maximum load, as well as increasing creep load. Anelasticity is thought to be caused primarily by the relaxation of subgrain boundaries bowing around pinning precipitates.  occurred in a ductile manner for all specimens tested, fracture characteristics include creep void formation and coalescence at the fracture surface. A high density of creep voids was found at grain boundaries and precipitates. Failure characteristics were more comparable to those reported in monotonic creep studies than performed in pure-fatigue or strain-controlled creep-fatigue studies found in existing literature.  At testing temperatures of 600 and 650 C, the same mechanisms are operating during load-controlled CF and monotonic creep, producing similar stress exponents and microstructural response to loading. However, subgrain coarsening occurs more quickly under CF loading conditions, resulting in higher strain rates and shorter rupture lives. It is postulated that the repeated backward and forward subgrain bowing that occurs during loading, unloading and at minimum load reduces substructure stability and accelerates subgrain coarsening, leading to increased strain rates and shorter rupture lives when compared to monotonic creep experiments.

Weldments of Grade 91 steel:  Cold metal transfer welding creates weldments with low heat-inputs compared to other arc welding processes. Welds from this study were produced using a heat-input of 0.67 kJ/mm.  CMT specimens exhibit an order of magnitude lower minimum strain rate than FCAW specimens under the same load-controlled CF conditions. The time-to-rupture is at least 5x longer for the CMT specimens in all test conditions examined. However, failure occurs in the ICHAZ for all specimens, regardless of welding process, which is indicative of type IV failure.  The CMT low heat-input welding process produces a HAZ which is 40% narrower than the FCAW HAZ. The cooling rate in the ICHAZ for the CMT specimens is 3x faster than the FCAW process.

 Previous studies suggest that M23C6 precipitates are critical for retaining elevated temperature strength in CSEFS including G91. The particle size is almost identical for CMT and FCAW after PWHT, however, the area fraction of M23C6 is higher in the CMT

6 specimens. After CF failure, the M23C6 particles in the ICHAZ of the FCAW specimens show more coarsening, and larger interparticle spacing when compared to those in the CMT specimens. Correspondingly, CMT weldments also exhibit greater substructure stability than FCAW weldments.  Simulated CMT and FCAW ICHAZ bulk specimens are also tested. Results from the simulated specimens are comparable to the FCAW full weldments, with only the CMT full weldments exhibiting enhanced CF strength.  DICTRA simulations show that carbon diffusion from the FZ to the ICHAZ may be a main factor contributing to the increase in area fraction of M23C6 and refined particle distribution in CMT after PWHT and CF deformation. The CMT weldment’s carbon enriched metal and narrow HAZ are both necessary to produce the favorable M23C6 distribution.

Uncovering the mechanisms governing creep-dominated creep-fatigue deformation and damage in Grade 91 base metal as well as type IV failure in Grade 91 welds contributes new insights that may lead to more accurate life prediction models than those currently in use. Moreover, the superior CF properties obtained by low heat input CMT has a high potential to develop practical solutions to the premature failure of Grade 91 welds in the current and future generation of nuclear power plants.

b. Introduction and Objectives Creep-strength enhanced ferritic steels (CSEFS) constitute a key class of structural materials due to their excellent creep-resistant properties. These alloys are currently being utilized in fossil-fueled power generation and petroleum processing, and are also being considered as candidate materials for applications in the next generation of nuclear fueled power generation, such as the sodium-cooled fast reactor.[1]–[3] CSEFS are low-carbon steels and were originally developed in the 1940s, the major alloying elements being 2 ¼ wt% chromium and 1 wt% molybdenum.[1], [4] Cr and Mo remain the principal alloying elements today. Figure 1 illustrates the evolution of CSEFS and corresponding maximum operating stresses for the major Cr-Mo CSEFS developed over the last 80 years, starting with the original 2 ¼Cr -1Mo CSEFS developed in the 1940s to the impressive rupture strengths of the latest generation of MARBN CSEFS which are still being developed for use in future power generation applications.[5] These steels typically contain a tempered-martensitic matrix, which is precipitation hardened to increase elevated temperature strength. In service, CSEFS components are subjected to elevated temperatures and stresses, which can be cyclic in nature, leading to the accumulation of creep and creep-fatigue deformation. Typically, CSEFS components are welded, common welding processes for CSEFS include gas metal arc welding (GMAW), gas tungsten arc welding (GTAW), shielded metal arc welding (SMAW), or submerged arc welding (SAW).[2], [5]–[8] Creep and creep-fatigue failures typically

7 occur in the welded region, either in the weld metal or heat-affected zone (HAZ). Failures in the HAZ are known as type IV failure, usually occurring in the outer region of the HAZ, especially in the fine-grained or inter-critical heat-affected zone (FGHAZ/ICHAZ). Significant debate continues in the welding literature concerning the mechanisms governing type IV cracking during creep and creep-fatigue deformation. Many studies indicate that the loss of elevated temperature strength in CSEFS is related to the microstructural changes that occur in the HAZ during welding. Regions where failure occurs typically reach temperatures between A1 and A3. Some have indicated that changes in precipitate distribution in this temperature range lead to loss of strength in the FGHAZ/ICHAZ.[7], [9], [10] Other studies indicate that accelerated creep damage in the outer region of the HAZ is related to the formation of ferrite and other microstructurally weak phases during cooling from welding.[11], [12] In addition, theories exist which attribute accelerated failures to the formation of stress triaxiality in the outer region of HAZ which increases strain accumulation in the outer region of the HAZ.[13], [14] From the multitude of mechanisms proposed to explain type IV failure in CSEFS welded components, it is clear that further exploration is necessary to explain this phenomenon especially with respect to creep-fatigue failure. While the majority of elevated temperature deformation studies for CSEFS are conducted under monotonic creep conditions, in service these components are usually subjected to cyclic stress at elevated temperatures, due to fluctuating loads and frequent component shut downs. [15] The cyclic nature of loading for CSEFS leads to the accumulation of creep-fatigue damage, which has been identified as a primary concern for future CSEFS development by power generation authorities. [16] ASME models for CSEFS life-prediction are currently developed using pure creep and pure fatigue mechanical data, collected in an uncoupled manner, which may not accurately reflect coupled creep-fatigue damage. [16], [17] Mechanistic knowledge incorporating microstructural evolution under creep-fatigue deformation is necessary to understand this type of damage and create better deformation models for CSEFS components subjected to CF deformation. This body of work begins with a review of the current state of the field and knowledge gaps for elevated temperature deformation of CSEFS base metal and weldments. Section c reviews the existing literature of the monotonic creep properties of V-modified 9-12 wt% chromium CSEFS base metal and weldments. In addition, the current state of research in creep- fatigue deformation of CSEFS is examined, investigating knowledge gaps and opportunities for exploration. This section identifies the many different mechanisms proposed for premature failure in CSEFS base metal and weldments under CF conditions, as well as the future of CSEFS, and current alloy development that could improve CSEFS performance and ultimately increase power generation efficiency. Section d describes the experimental procedures used in the present study to examine microstructural evolution after welding and deformation of Grade 91. The primary type of mechanical testing performed in the current study is creep-fatigue testing. Section d explains the experimental setup for creep-fatigue testing. Several types of characterization techniques

8 are utilized to fully explore the microstructural response to processing and deformation of the specimens examined in this body of work. The primary mode of characterization for this study is electron microscopy. Each characterization technique is explained in detail, including sample preparation. Also, analytical techniques utilized in the current study such as electron backscatter diffraction (EBSD) and energy dispersive spectroscopy (EDS) are described. Section e investigates base metal creep-fatigue properties of the G91 alloy. The objectives of the base metal studies are to assess the effects of loading parameters and temperature on mechanical response to load-controlled, creep-dominated CF deformation in CSEFS. In addition, microstructural evolution during CF deformation is investigated, focusing on microstructural features that are important to elevated temperature deformation in CSEFS, including defect content and precipitate distribution. Significant amounts of anelastic backflow are observed during load-controlled CF deformation of G91 base metal. This anelasticity is explored and the consequences of anelastic motion of subgrain boundaries on CF deformation are examined. Theories to explain the significant decrease in rupture life for CSEFS under CF conditions compared to monotonic creep are hypothesized. Section f explores the creep-fatigue properties of G91 weldments. The microstructural changes that occur in the ICHAZ, where failure occurred for all weldment CF tests conducted, is examined. Most importantly, the effect of low-heat-input welding on ICHAZ microstructure and CF performance is assessed. It was determined that through a combination of appropriate filler metal composition and narrow HAZ widths achievable by low-heat input welding, a favorable precipitate distribution is achieved in the ICHAZ, leading to significantly increased ruptured lives compared to those achieved using conventional welding techniques and filler . The low- heat-input process utilized for the CF weldment studies is known as cold-metal-transfer (CMT), a recently developed, commercially available arc welding process, which could possibly be implemented in the field for welding CSEFS power generation components. Section g summarizes this body of work, the contribution to the field and highlights unanswered questions that require further investigation. Future work that could further the field of G91 and other CSEFS, and ultimately lead to enhanced power generation capabilities for the ever-increasing energy demands of today’s civilization are discussed.

c. Background c.1 Creep-strength Enhanced Ferritic Steels The necessity for high-efficiency energy production has increased demand for structural materials used at elevated temperatures for power generation. An obvious contender to meet these structural requirements resides in a class of alloys known as creep-strength enhanced ferritic steels (CSEFS). CSEFS make excellent elevated temperature structural materials for fossil-fueled and nuclear power applications due to mainly their creep resistant properties These economical alloys also exhibit low thermal expansion, and low-activation, which make them suitable for applications where welding is necessary, and radiation may be present. In

9 addition, depending on alloy chemistry, CSEFS can exhibit excellent oxidation resistance as well.[3]–[5] CSEFS have been utilized in high-temperature structural applications for over half a century due to their superior creep resistance.[4], [18] Depending on alloy composition, heat treatment, and mechanical processing, microstructures can be ferritic, martensitic, bainitic, or a mixture of the aforementioned. Typical microstructures include prior austenite grains (PAGs) within which exists martensite or bainite. At finer length scales, CSEFS contain a high dislocation density, a fine subgrain network and a fine distribution of second phases in the normalized and tempered condition. An overview of common alloying elements for CSEFS has been summarized effectively in the review by Klueh. [3] A more concise summary of the major alloying elements is found below: Carbon and Nitrogen: Added to stabilize austenite and form martensite upon cooling. Carbon and Nitrogen are potent precipitation hardeners via the formation of carbides upon tempering after normalization. Chromium: Added for oxidation and resistance, Cr also forms carbides which can increase elevated temperature strength depending on precipitate type and distribution. Molybdenum and Tungsten: Added as a solid solution strengthener. Mo can also be consumed by second phases depending on alloy content and processing. Common second phases rich in Mo and W are Laves phase, M6C and M2C.[19] Mo-rich carbides can be advantageous or detrimental to creep resistance depending on precipitate type, size and distribution. On the other hand, the formation Laves phase is typically thought to be detrimental if the alloy also contains Tungsten.[20] Vanadium, and Niobium: Vanadium and Niobium are strong carbonitride formers, causing the formation of fine MX particles which are effective precipitation hardeners. The addition of vanadium led to the development of Modified G91 which produced a marked increase in creep rupture strength, attributed to the formation of a fine distribution of VX carbonitrides. [8] Boron: Boron has been recently added to the latest generations of CSEFS. Additions of boron with careful control of nitrogen content have been shown to significantly increase creep strength of CSEFS. However, contention exists regarding the mechanisms behind the increase in creep strength.[21]–[23] Also, addition of Boron and Nitrogen has been found to suppress fracture in the FG/ICHAZ which is advantageous for creep strength of CSEFS welded components.[24], [25] Manganese, and Silicon: Mn and Si are both austenite stabilizes, making it easier to form martensite upon cooling for CSEFS. These elements also act as oxygen and phosphorus getters, removing O and P from solution, which can otherwise be detrimental to elevated temperature strength.

10 CSEFS were first introduced in the 1940s for boiler applications with 2.25Cr-1Mo alloys, and these materials were developed as an alternative to austenitic stainless steels, which could handle higher operating temperatures but suffered from large coefficients of thermal expansion leading to significant thermal fatigue during plant shutdowns. In the 1970s, improvement in alloying resulted in the development of German grade X20 (12Cr-1-Mo-0.25V) and G91 (9Cr-1Mo-0.25V); the latterwas developed at Oak Ridge National Lab in the United States. After the development of G91, CSEFS were also utilized in petro-chemical applications in petro-chemical hydrocracking furnaces.[2], [4] Table 1 shows the typical compositions for the current and previous generations of CSEFS used in elevated temperature applications.[3]

Table 1: Composition in wt% for current and previous generations of Cr-Mo CSEFS.[3]

c.2 Alloy Development for Elevated Temperature Environments The 2.25Cr-1Mo steels developed in the 1940s have creep-rupture lives over 100,000 hours at 35 MPa and 600°C.[26] G91 has a 100,000 hr creep-rupture strength of 100 MPa at 600°C.[26] The nominal composition of G91 can be found in Table 2. While G91 was not the first 9Cr-1Mo steel developed, additions of niobium and vanadium as well as a nitrogen specification resulted in notable creep-strength increases over previous generations of 9Cr ferritic-martensitic alloys. [27]

Table 2: Nominal composition of G91 steel in wt. %.

Element C Cr Mo V Ni Mn N Nb B Si Al Cu

Composition 0.08 8.40 0.92 0.23 0.09 0.45 0.04 0.075 10-4 0.36 0.02 0.08 wt%

11

G91 is usually employed in the normalized and tempered condition, which involves normalizing at 1060°C followed by air-cooling to produce a fully martensitic microstructure. The martensite is then tempered at 760°C.[28] The resulting microstructure consists of low- angle subgrain boundaries (LAGB), high-angle prior-austenite grain boundaries (PAGB) as well as high-angle block and packet boundaries.[29] The subgrain interior has significant dislocation 14 -2 content prior to deformation (ρdis >10 m ).[30] Second phase M23C6 particles precipitate at PAGB and subgrain boundaries. MX type particles precipitate within the subgrain interior.[27] Figure 2 illustrates the microstructure of G91 in the normalized and tempered condition.[31], [32] The primary phases, their space group, primary solute, and lattice parameters are shown in Table 3.

Table 3: Descriptions of matrix and secondary phases found in G91.

Lattice Parameter Phase Primary Solute Space Group (푨̇ )

BCC (tempered Fe, Cr, Mo 퐼푚3̅푚 2.867 martensite)

M23C6 (Cr, Mo)23(C,N)6 퐹푚3̅푚 10.687

MX (V,Nb,Mo),(C,N) 퐹푚3̅푚 4.165

9-12% Cr CSEFS benefit from several high-temperature strengthening mechanisms.[33] The high dislocation density results in dislocation-dislocation interactions and reportedly provides a strain-hardening effect.[16] Chromium is added to provide corrosion resistance and also forms M23C6 particles which nucleate at PAGB and subgrain boundaries. M23C6 precipitates are thought to provide a boundary pinning effect, preventing subgrain coarsening and adding elevated temperature strength. M23C6 particles are also thought to inhibit mobile dislocation motion by obstructing dislocation knitting reactions at LAGBs.[27], [34] Vanadium and niobium rich second phases precipitate in the form of MX type carbides and carbonitrides. MX precipitates nucleate within subgrain interiors and impede dislocation motion during creep.[31] Many studies have examined the creep properties of these alloys. However, questions remain concerning the relative contributions of each strengthening mechanism to the overall creep strength of CSEFS, especially regarding the effects of second phase precipitation on elevated temperature strength. Kostka et al. [29] argued that the main source of creep strength in CSEFS is a substructure pinning effect by M23C6 type carbides. They tested this theory by comparing creep results from a 10% chromium CSEFS to a Fe-10Cr alloy which had been processed to produce similar grain size and dislocation density, without second phase precipitation. Creep tests

12 revealed the strain rate from CSEFS creep specimens was several orders of magnitude lower than the Fe-10Cr alloy. While this experiment emphasizes the significance of precipitation strengthening, unfortunately, it does not isolate the strengthening effect of the M23C6 precipitates.

The Fe-10Cr alloy did not contain boundary strengthening M23C6 particles or fine MX type precipitates which contribute to strength via dislocation pinning, effectively removing both precipitation-strengthening mechanisms present in CSEFS, making it impossible to assess the contribution of subgrain boundary pinning to the overall creep resistance of the alloy. Kostka et al. [29] concluded that the high dislocation density present in CSEFS did not contribute to strengthening, because the dislocation densities of the Fe-10Cr alloy and the CSEFS were similar before creep-deformation, and because dislocation density in CSEFS decreases dramatically at the onset of creep deformation.[29] While the dislocation density does decrease, the dislocation density after severe creep deformation is still significantly large in CSEFS, as demonstrated in experiments conducted by Pesicka et al. [30], who found that dislocation densities decrease from 9.24 x 1014 m-2 to 0.29 x 1014 m-2 after accumulating 8% creep strain in X20 (12Cr-1MoV). Kostka et al. [29] stated that dislocation density in both alloys was similar before creep but did not report dislocation densities after creep. It should be noted that the difference in boundary character between the two alloys (low-angle subgrain boundaries in the CSEFS vs. high angle grain boundaries in the Fe-10Cr alloy), was not considered. Cerri et al. [32] performed short-term creep tests to examine the effect of creep deformation on the microstructure of G91 steels. They argued that MX type precipitates are not present in the initial microstructure of G91 steels and precipitate only after the onset of creep deformation, concluding that the only precipitate present after tempering is M23C6. These observations were based on precipitate size measurements, which followed a Gaussian distribution with a single peak prior to creep and a bimodal distribution after the onset of creep. The average size of the precipitates prior to creep deformation was 79 nm. The type of precipitate could have been confirmed based on precipitation location, precipitate composition using energy dispersive X-ray spectroscopy (EDS), or diffraction experiments. It should be noted that the tempering treatment was only one hour, which may not have been sufficient time to nucleate MX type particles of resolvable dimensions. Other studies on CSEFS report MX type carbides present after tempering, however, most experimenters employ tempering treatments of at least 2 hours at 760C.[20], [21], [32], [35], [36] The literature studies summarized above demonstrate the knowledge gap concerning the contributions of the various elevated temperature strengthening mechanisms present in CSEFS as well as the evolution of the microstructure during elevated temperature deformation. A better understanding concerning the effects of various strengthening mechanisms as well as microstructural development during creep is necessary to formulate accurate deformation models for these alloys and improve future generations of CSEFS.

13 c.3 Creep-fatigue: A More Accurate Description of Deformation in G91 Components Traditionally, creep performance has been the major design criteria for high- temperature structural materials used in the energy sector, resulting in a wealth of knowledge concerning the creep properties of G91 and other CSEFS. Recent studies indicate a more accurate description of mechanical deformation occurring in high-temperature structural materials for power generation is a combination of creep and fatigue damage, termed creep- fatigue (CF). In power plants, CF deformation arises due to pressure fluctuations during operation as well as frequent component shut downs.[15] In a status report conducted by Argonne National Laboratory, CF deformation was recognized as the most critical issue in high- temperature structural design for power applications, with emphases on the lack of creep- fatigue data and mechanistic knowledge for G91 steels and weldments.[16], [17] Due to the lack of available CF data, ASME adopted life-prediction models (known as “interaction diagrams”) which combine creep and fatigue data collected at laboratory time scales.[16] The majority of CF tests are conducted using strain-controlled testing conditions. The most obvious mechanical deformation feature present in all CF experiments conducted using strain-controlled testing devices is the pronounced cyclic softening exhibited by G91 specimens. Concerning the materials approved by ASME for high-temperature use in nuclear power components, cyclic softening is unique to CSEFS. Austenitic stainless steels, for which the ASME life-prediction models were originally designed, cyclically harden when exposed to CF environments.[16] Examples of cyclic softening in G91 CF specimens are shown in Figure 3. Kim and Weertman showed that addition of a two-minute hold time at maximum strain results in a decrease in cycles to failure for G91 specimens, compared to pure fatigue.[37] In 2008, Takahashi investigated the effect of hold time further, performing over 80 creep and CF tests with dwell times varying from 0 to 10 hours. From these experiments it is evident that dwell time has a significant effect on rupture life of CSEFS, as shown in Figure 4. [15] Kim and Weertman, Takahashi, and Gopinath et al. all reported a substantial reduction in lifetime when a dwell-period is introduced in compression compared to tension hold. [15], [37], [38] Results from one study indicate a dwell time of two minutes in compression results in a life-reduction of 40% compared to two minutes in tension. [37] However, as shown in Figure 4, at longer dwell times the cycles to failure approach similar values for tensile and compressive stress states. [15] This experimental evidence illustrates a limitation associated with the current creep-fatigue database available for life-prediction models. The majority of creep-fatigue data available is for relatively short dwell-times, which is not necessarily appropriate for predicting the extended dwell-times associated with plant operation. The decrease in rupture life associated with compression dwell is sometimes attributed to a combination of CF deformation and environmental .[39]–[43] CF experiments on CSEFS conducted under vacuum indicate that dwell time under tensile stress results in shorter lifetimes than dwell time in compression. [37] Few studies examine the difference in fine-scale microstructural evolution in tension CF compared to compression CF. Qualitative experimental evidence by Gopinath et al. suggests

14 that substructure coarsening occurs more rapidly during compressive holds than tensile holds, indicating the possibility of not only an environmental effect on crack propagation but also a difference in the microstructural evolution under each type of deformation. [38] Fournier et al. discovered through a series of load-controlled and strain-controlled creep-fatigue experiments on G91 that load-controlled deformation results in significantly larger viscoplastic strain accumulation during peak load, which causes a reduction in rupture- life for G91 CF specimens. This finding raises concerns over whether or not it is appropriate to formulate life-prediction models based on strain-controlled CF data.[39]–[41] While the link between microstructural evolution and monotonic creep deformation has been explored extensively for CSEFS, there have been significantly fewer studies concerning microstructural evolution under CF deformation conditions. The majority of CSEFS CF studies to date emphasize CF data collection and the formulation of mechanical data-based life-prediction models. Kim and Weertman made several qualitative observations using TEM and concluded that cyclic-softening observed during strain-controlled CF experiments was due to subgrain and carbide coarsening. [37] Gopinath et al. observed integration of the initially high dislocation density into LAGB in P92 steels and determined that this was the dominant softening mechanism during elevated temperature deformation. While Gopinath et al. did not explicitly mention subgrain coarsening as a deformation mechanism, it is likely that the reduction of dislocation density and subgrain coarsening are interrelated, and both play a role in the cyclic softening behavior of CSEFS. They also witnessed significant dislocation pinning by fine MX carbides and nitrides present in subgrain interiors, shown in Figure 5. [38] This finding is significant, especially when compared to the findings of Kostka et al. [29] who argued that Orowan type dislocation pinning interactions rarely occur in CSEFS. There have been few attempts to quantitatively correlate the microstructural evolution of CSEFS with CF deformation. [38] For the design of safe and efficient reactors and fossil fueled steam turbines, development of accurate life-prediction models for CF deformation in CSEFS is of critical importance. One common life-prediction model is the time-fraction approach (TFA) currently adopted as the design criteria for G91 components in nuclear applications by ASME. The TFA is based on empirical, uncoupled creep and fatigue data collected at lab time scales and extrapolated to service lifetimes. [16] Creep-damage is represented by Equation (1) below and the vertical-axis in Figure 6. In this equation, dc is the creep damage per cycle, tH (s) is the dwell period of the CF life being predicted,  (MPa) is the stress applied, Tabs (K) is the test temperature and tR (s) is the rupture time of a monotonic creep test at the same stress level and temperature of the CF test in question. Fatigue damage is represented by Equation (2) and is plotted on the horizontal-axis of the CF interaction diagram in Figure 6. In Equation (2), df is the fatigue damage per cycle, Nf is the cycles to failure for a pure fatigue test with the same strain range, ∆휀, is the strain range, and 휀̇ is the strain rate (1/s). Fatigue damage is the reciprocal of the pure-fatigue life for a fatigue test of similar strain range, strain rate, and temperature as the CF test being predicted. ASME has chosen the intersection point for the

15 TFA to be (0.1,0.01), if failure occurs within the failure envelope shown in Figure 6a, the component is not adequate for service. [15], [16], [42], [44]

푡퐻 푑푡 푑푐 = ∫ (1) 0 푡푅(휎,푇푎푏푠)

1 푑푓 = (2) 푁푓0(∆휀,휀̇ ,푇푎푏푠)

푡퐻 휀̇ 푑푐 = ∫ 푑푡 (3) 0 훿(휀̇ ,푇푎푏푠)

The other major life-prediction model employed for CSEFS is the ductility exhaustion model (DEM), which differs from the TFA in the way creep damage is calculated. For the DEM approach, inelastic strain-rate and creep ductility (훿) are used to calculate creep damage (Equation (3)), as opposed to the stress-based time-to-rupture function in Equation (1). [15], [42] The failure predictions shown in Figure 6 were calculated based on tests conducted by Takahashi. The TFA failure envelope determined by ASME is also shown. [15], [45] Figure 6a shows that the majority of rupture lives predicted using the TFA lie outside the ASME envelope while all of the predicted rupture lives occurred well outside the failure envelope when using the DEM approach shown in Figure 6b, illustrating the over conservatism associated with the failure envelope determined by ASME. Figure 7 shows a comparison of the life prediction results from Figure 6 to experimental rupture lives from CF experiments conducted by Takahashi [15]. Figure 7a shows the slightly over-conservative life predictions and large scatter associated with the forecast using the TFA. Figure 7b shows the vastly over-conservative life predictions calculated using the DEM. Takahashi concluded that the over-conservatism associated with the DEM was the result of an overestimation of creep-damage (as evidenced by Figure 6b). Takahashi argued the accelerated damage-rate during the beginning of a strain hold was more damaging than the later portions of the strain hold, which had a lower damage rate. A modified DEM was proposed which subtracts the ductility not due to creep-damage (calculated from elevated temperature tensile testing) from the ductility contributions that were due to creep damage. The modified equation for creep damage is shown in Equation (4), where 훿0 is the non-creep ductility. [15]

푡퐻 휀̇ 휀̇ 푑푐 = ∫ − 푑푡 (4) 0 훿(휀̇ ,푇푎푏푠) 훿0(푇푎푏푠)

16 The result of the modified ductility equation is shown in Figure 7c. The scatter between predicted and actual failure falls between a factor of two for all experiments performed. Similar attempts to improve the DEM model have been performed by Aoto and Fournier et al. [39]–[42] One particularly important aspect of the CF process that is often “muted” by the use of strain-controlled testing is the anelastic strain which causes backflow during each cycle in load- controlled testing. Due to the nature of displacement-controlled testing, backflow of strain is not possible at minimum load. Schematic load-controlled and strain-controlled CF hysteresis are shown in Figure 8.[39] Even the Fournier’s experiments which utilized load-controlled waveforms did not allow for anelastic strain at minimum load, due to the fact that loading was switched to strain-control for all portions of the hysteresis except for the portion at maximum load. Anelastic backflow of strain in metals that form substructure was observed as early as the 1980s, when Gibeling and Nix determined that upon abrupt stress unloading during creep, class II alloys would experience a backflow of strain. Gibeling determined that the anelasticity was due in part to the bowing and un-bowing of subgrain boundaries during loading and unloading. [46], [47] CSEFS are similar to class II alloys in the respect that they both exhibit substructure formation. However, a significant difference lies in the fact that CSEFS start with a refined substructure prior to deformation (a consequence of the martensitic transformation) and class II alloys start with no substructure prior to creep damage. The substructure in class II alloys develops as a consequence of creep damage and the rearrangement of into low energy structures. Sawada and Kimura realized this correlation and decided to study the effect of abrupt stress unloading on the creep behavior of 9 wt% Cr CSEFS. Sawada et al. discovered that upon abrupt stress unloading CSEFS did exhibit anelastic recoverable strain if the CSEFS exhibited a subgrain network. In addition, Sawada discovered that if the material contained a carbide stabilized subgrain network, the amount of anelastic backflow after abrupt stress unloading was decreased compared to a material with a non-carbide stabilized substructure. Figure 9 illustrates a comparison of the anelastic backflow resulting from a CSEFS with a carbide stabilized subgrain network and a CSEFS with a non-stabilized subgrain network, as well as a CSEFS with a fully ferritic microstructure and no subgrain network. [48] Sawada also found that abrupt stress unloading only resulted in anelastic backflow at elevated temperatures, indicating the anelasticity is a thermally activated process. [48] Figure 9c shows the displacement curve for a ferritic steel with no lath structure, indicating that without a subgrain network, no anelastic recovery will occur after an instantaneous stress reduction. Figure 9a shows P92, a CSEFS with a substructure stabilized by M23C6 and Laves phase. [35] Figure 9d shows a steel with a martensitic substructure that is stabilized by a fine distribution of MX carbonitrides on lath boundaries. A comparison of the displacement curves in Figure 9 shows the following two main points: first, substructure must be present for a CSEFS to exhibit anelastic recovery, and second, the amount of anelasticity

17 observed will be heavily dependent on how well the substructure is stabilized. Sawada also reported that the steel with the fine distribution of MX carbonitrides at lath boundaries significantly outperformed P92 under creep conditions as well, suggesting a correlation between resistance to anelastic recovery and resistance to creep deformation. [48] The anelastic backflow phenomenon described above should play a significant role in load-controlled CF deformation, due to the fact that it occurs during every load-controlled CF cycle, as evidenced by the schematic load-controlled CF stress-strain hysteresis shown in Figure 8. To fully understand the role of anelastic backflow on CF deformation, a systematic study is necessary examining the effect of CF loading parameters on anelastic backflow as well as the correlation between anelastic backflow and forward flow of creep-fatigue strain. c.4 Weldability Issues Associated with 9-12% Cr CSEFS While the performance improvements associated with CSEFS over the last four decades have been promising, their full operating potential is restricted by frequent premature failures in welded sections of CSEFS components. [5] CSEFS can be welded using a variety of processes, the most common processes including: gas metal arc welding (GMAW), flux-cored arc welding (FCAW) and gas tungsten arc welding (GTAW). [6], [49], [50] Filler metal used for welding 9- 12% Cr CSEFS is typically designed to match the elevated temperature properties of the parent metal, most notably the creep strength. [18], [51] Premature failure in CSEFS can be attributed to the microstructural gradient produced during welding in the parent material adjacent to the fusion zone, known as the heat-affected zone (HAZ).[6], [52] The microstructural gradient in the HAZ has been summarized effectively by Mannan and Laha, and is illustrated schematically in Figure 10. [18], [53] The Coarse-Grained HAZ (CGHAZ) reaches a temperature above Ac3 during welding and all carbides present prior to welding are dissolved, resulting in large PAGs after cooling. The fine-grained heat-affected zone (FGHAZ) reaches a temperature above Ac3 and carbides present prior to welding are partially dissolved. Remaining carbides limit the growth of austenite grains during cooling. The temperature in the inter-critical HAZ (ICHAZ) is between Ac1 and Ac3, resulting in partial formation of new austenite and subsequently fresh martensite as well as over-tempering of the tempered martensite in the parent material. It is widely accepted that premature failure in G91 and other CSEFS components is associated with the HAZ, particularly the FGHAZ or ICHAZ, where failure usually occurs, termed type IV failure. [5], [10], [13], [18], [39], [52], [54] While many different mechanisms have been proposed regarding type IV failure in G91 and other CSEFS, contention exists regarding the accuracy of each. Many studies suggest the loss of creep-strength in G91 weldments is due to changes in precipitation behavior as a result of welding. [7], [9], [10], [55] Hirata and Ogawa performed TEM on simulated HAZ specimens and base metal to determine precipitate evolution during welding. They postulated that partial dissolution of M23C6 precipitates during welding causes an enrichment of chromium in the matrix. The chromium enrichment results in precipitation of

18 fine M23C6 precipitates surrounding the partially dissolved M23C6 precipitates. The distribution of fine M23C6 precipitates surrounding larger M23C6 precipitates leads to Ostwald ripening during creep and an increased coarsening rate in the FGHAZ, resulting in accelerated creep deformation in the FGHAZ. They also predicted that chromium enrichment in the matrix would promote coarsening of MX type precipitates, however, experimental evidence suggests MX precipitate growth is controlled by the diffusion of vanadium, not chromium. [10] Yu et al. suggested the loss in creep-rupture strength in the HAZ was due to changes in precipitate distribution during welding. Pre-tempering the base-metal at a lower temperature prior to welding achieved a finer precipitate distribution. The lower pre-heat temperature resulted in a creep-rupture life 5x greater than specimens pre-heated at the standard pre-heat temperature. [7], [49] The increase in rupture life was attributed to a finer distribution of M23C6 carbides present in the weldment with a lower pre-heat temperature. Yu calculated Orowan bowing stresses for M23C6 to be 71 MPa and 51 MPa for the low-temperature and high- temperature pre-heating conditions respectively. Based on these calculations, Yu concluded that the fine distribution of M23C6 precipitates increased resistance to dislocation motion. However, experimental evidence from other studies suggests that dislocations interact mainly with MX type precipitates, raising concerns over the validity of describing the strengthening effects of M23C6 precipitates using an Orowan type mechanism, usually reserved for precipitation strengthening via dislocation-precipitate interactions. [38], [56] Quantification of the M23C6 strengthening effect may be achievable via a Zener pinning type mechanism, since M23C6 precipitates are thought to pin subgrain boundaries during creep. [29], [34], [57], [58] The number density of M23C6 carbides in the low-temperature pre-heat condition was over 3x larger than the number density of the high-temperature heat treatment. The increased creep strength may originate from increased resistance to subgrain coarsening from a greater density of M23C6 carbides at PAGB and subgrain boundaries. In subsequent studies, Yu et al. attributed the reduced creep strength in G91 weldments to the formation of ferrite in the FGHAZ.[12] Yu suggested that increased chromium content around partially dissolved M23C6 carbides promotes formation of ferrite rather than martensite when cooling from austenite. They used electron backscattered diffraction (EBSD) image quality (IQ) maps to show fresh ferrite formation after tempering. While IQ maps are sensitive to strain in the lattice, they are also sensitive to other factors such as surface relief. [59], [60] Another way to use EBSD to find fresh ferrite could be to utilize kernel average misorientation (KAM) or grain reference orientation deviation (GROD) maps which show orientation differences between adjacent areas in a grain. Examples of IPF and IQ maps which Yu argued showed fresh ferrite are shown below in Figure 11. Local misorientation within a fresh ferrite grain should be significantly lower than the tempered martensite structure, due to the differences in dislocation density. GROD maps were shown to have the best correlation to STEM dislocation density measurements performed in previous studies on Ni-based superalloys. [60] Wang et al. also reported similar observations of ferrite formation in G91. [11]

19 Other type IV failure mechanisms have been associated with the microstructural gradient present in the weldment as opposed to changes in microstructure occurring specifically in the FGHAZ. Many attempts have been made to examine elevated temperature deformation in the weldment microstructure through modeling efforts. [9], [13], [61], [62] Li et al. attempted to model creep-deformation in a CSEFS (P122) by testing simulated HAZ samples along with base metal (BM) and full weldment (WM) samples and applying their individual mechanical properties to a finite element model (FEM). They used a classic Norton law to model the primary, secondary and tertiary regions of the weldment creep test, resulting in the simulated creep curve shown in Figure 12. [13] A more accurate model may have been produced by only modeling the secondary or “steady state” creep regime or by modeling the primary creep-regime using a parabolic creep relationship. A FEM model was developed using mechanical test data from simulated specimens (Figure 12b) with mechanical property values assigned to each portion of the weldment. The size of each section (BM, FGHAZ, CGHAZ, etc) was determined by taking hardness profiles across the weldment. While the mechanical properties in an actual weldment vary continuously with microstructure, each section in the FEM model was discontinuous, which could lead to some error associated with the model’s results. Li found that the triaxial stress state and equivalent strain was significantly higher in the FGHAZ than other sections of the weld. Li hypothesized that constraints in the FGHAZ from the BM and CGHAZ create large stress triaxiality in the FGHAZ which promotes creep-void formation and accelerates failure in the FGHAZ region compared to the rest of the weldment. This work is questioned by findings from Albert et al. who showed that simulated FGHAZ specimens failed before weldment specimens (Figure 13a), indicating that the premature failure may be a function of the microstructure in the FGHAZ as opposed to the microstructural and mechanical property gradient produced in the HAZ. [6], [55] Results from models produced by Li et al indicate that reducing HAZ width also reduces the triaxial stress state and therefore increases rupture life. Albert et al. confirmed that changing the HAZ width changes creep strength by performing experiments using welding processes with different heat inputs, which produced varying HAZ widths, shown in Figure 13b. [6] c.5 The Future of CSEFS Recent advancements in alloying have led to promising improvements in the elevated temperature deformation properties of ferritic- martensitic steels. The next generation of these steels (designated P92) has primary alloy additions of tungsten (approximately 1.5%) as well as a reduction in molybdenum. [3] Tungsten acts primarily as a solid solution strengthener but also precipitates out of solution to form Laves phase. The microstructure present in P92 after tempering strongly resembles the microstructure of G91, with similar dislocation density, subgrain size and precipitate distribution. [3], [35], [36], [63] A comparison of creep-rupture strength for G91 and P92 is found in Figure 14a. [24]

20 Sawada et al. hypothesized that precipitation of fine Laves phase occurs at prior austenite and subgrain boundaries in CSEFS containing tungsten, resulting in a boundary pinning effect and increased rupture times compared to G91 as shown in Figure 14a. [63] While this would explain the decrease subgrain coarsening rates found in P92 (shown in Figure 14b), it does not explain the significant increase in dislocation density stability in P92 compared to G91 during creep, also shown in Figure 14b. It should be mentioned that the PAG size of the G91 and W containing CSEFS used in this experiment were 33μm and 20μm respectively. The PAG size may affect the creep properties of each material. Plesiutschnig et al. reported that increasing the PAG size leads to an increase in creep rupture time for multiple CSEFS alloys. [64] A more accurate comparison may have been conducted by austenitizing both materials to have similar grain PAG sizes prior to quenching and tempering. Contrary to Sawada’s findings, Ennis et al. argued that the primary strengthening mechanism resulting from W additions was a solid solution effect, which prevented dislocation movement in subgrain interiors as well as the migration of subgrain boundaries. Ennis and Taneiki found that at temperatures above 600°C, accelerated coarsening of Laves phase became detrimental to the creep strength by removing W from solution and did little to strengthen the material. After 22000 hours of creep at 600°C Ennis reported dislocation densities in P92 2x larger those in G91 after 10000 hours of creep under the same conditions. The decreased rate of dislocation annihilation reported by Sawada et al supports Ennis’ claim concerning the solid solution strengthening properties of tungsten in CSEFS. Taneiki reported Laves precipitate sizes between 500-700 nm after creep at 650°C, which is approximately the same size as the subgrains, indicating a low probability for any effect on subgrain coarsening or dislocation interactions. [35], [65] Wang et al showed that welding P92 resulted in increased rates of Laves phase coarsening in the FGHAZ compared to BM, which Wang attributed to type IV failure during creep in P92 steels. [66] While the creep properties of P92 at lower temperatures have been promising, their mechanical properties at temperatures higher than 600°C as well as their poor weldability create concerns regarding the ability of P92 to enhance operating conditions of the power plants they are utilized in. Another promising development resulted from controlled additions of boron and nitrogen to CSEFS (designated MARBN steels). [24], [25], [64], [67]–[69] Abe, Albu, and Mayr et al. found that by adding controlled additions of B and N to CSEFS an impressive increase in creep strength could be realized, shown in Figure 14a. Albu and Abe found via Electron Energy Loss Spectroscopy (EELS) that boron replaced carbon in the chromium rich M23C6 carbides that nucleate on subgrain boundaries. [67], [69] Boron was found to reduce the coarsening rate of M23C6 precipitates, resulting in a reduction in creep rate and greater rupture life compared to G91 and P92 steels. [24] Abe performed a study on N content and found that N additions up to 79 ppm increased the creep strength of the MARBN steel while concentrations above this resulted in the formation of coarse BN precipitates that decrease creep strength by depleting the matrix of B and N while offering little precipitation hardening or boundary pinning effects. [67] No mechanisms were provided for the decreased coarsening rate or the effect of nitrogen content on creep life. One possibility for the increased creep rupture lives via nitrogen

21 additions could be that an increase in nitrogen allows for a greater number of MX nitrides to nucleate, which pin mobile dislocations during creep. Boron and nitrogen additions have also shown a propensity to suppress the FGHAZ formation in CSEFS weldments, greatly increasing the creep strength of CSEFS weldments, also shown in Figure 14a. Mayr, Abe and Liu studied the effects of FGHAZ suppression in B added CSEFS. [24], [68], [70] The mechanisms for FGHAZ suppression resulting from additions of B and N are still under consideration but are largely related to suppression of fine grains in HAZ. Investigations are also underway to study the effects of welding process on type IV failure and creep properties of CSEFS weldments. As previously stated, Yu found that changing the pre-weld tempering treatment in G91 to 650°C instead of the industry standard 760°C resulted in a creep-rupture life 5x greater than specimens that underwent the industry standard tempering treatment. [7], [12], [49] Albert et al found that using low heat-input welding processes resulted in finer HAZ widths and that optimizing groove angle could also result in increased rupture lives, shown in Figure 13b. Both of these process changes were attributed to a decrease in stress triaxiality in the HAZ similar to the results of Li referenced earlier. [6], [13] It should be noted that the low heat-input welds in Albert’s study were produced using electron beam and laser welding processes, both of which would be impractical for implementation in the field. However, a fairly new welding process known as cold metal transfer (CMT) can be implemented in a setup essentially the same as gas metal arc weld (GMAW) to produce low heat-input weldments. CMT has shown promise concerning low heat- input welding in other alloy systems. [71], [72] Finally, Albert et al showed that PWHT had little effect on the suppression of type IV cracking in weldments. [14]

c.6 9-12% Cr CSEFS: Summary and Unanswered Questions Significant progress has been made studying the elevated temperature properties of G91 and other 9-12% Cr CSEFS. The preceding literature review has emphasized a few of the remaining questions requiring further exploration. Kostka, Eggeler and Cerri et al tracked microstructural evolution and correlated the monotonic creep strength of CSEFS to their microstructure. [29], [32], [34], [73] However, as illustrated above, questions remain regarding the relative contributions of each strengthening mechanism to the overall creep strength of G91 and other CSEFS. Knowledge of the relative contributions of each strengthening mechanism is crucial for successful improvements concerning the mechanical properties of future CSEFS. Concerning CF interactions in CSEFS, significant progress has been made in the form of mechanical testing and empirical life prediction models based on CF data. [15], [38], [39], [42] While Gopinath et al, Kim and Weertman and others have made qualitative microstructural examinations, few attempts have been made to quantitatively examine microstructural evolution with CF strain and determine microstructural response to CF testing parameters such as dwell time, temperature and applied stress. [37], [38] Quantitative microstructural investigations regarding the CF properties of CSEFS may lead to microstructure- based life-prediction models for current and future generations of CSEFS.

22 Many different mechanisms have been suggested to explain the type IV cracking phenomenon present in CSEFS including: changes in precipitation behavior suggested by Hirata and Ogawa, formation of fresh ferrite in the FGHAZ during welding suggested by Yu, and Wang, or stress triaxiality in the HAZ, suggested by Li et al. [7], [10], [12], [54] While the exact mechanism causing type IV failure is still under consideration, promising advances have been made concerning improving the elevated temperature mechanical properties of CSEFS and weldments. The addition of boron and nitrogen to CSEFS has enhanced the creep properties compared to G91 and P92 and resulted in the elimination of the FGHAZ in CSEFS weldments. The mechanistic details concerning the creep strength and HAZ suppression in MARBN steels require further investigation. [24], [67], [69] Studies concerning welding processes used for CSEFS have shown that using lower pre-heat temperature, different weld geometries and low heat-input welding processes may mitigate the type IV cracking issue. However, whether or not low heat-input arc welding processes such as CMT can be successfully employed to weld CSEFS has not been determined. [6], [7] Knowledge of microstructural evolution with creep and CF deformation in CSEFS and weldments is crucial to obtain a better understanding of the deformation mechanisms operating in these alloys. Mechanistic knowledge concerning deformation in G91 and weldments may create paths for improved life-prediction models and performance improvements via adjustments in alloying, processing, and fabrication techniques.

d. Experimental Procedures d.1 Materials and Welding d.1.1 G91 Base Metal 19 mm thick G91 plates were received in the normalized and tempered state from American Alloy Steel. Normalization consists of an austenitization treatment at 1040C followed by air cooling. A subsequent one-hour tempering treatment at 760C followed by air cooling was performed. In service, the majority of CSEFS components are welded and post- weld heat treated (PWHT) to relieve internal stresses, temper fresh martensite and develop favorable precipitate distributions. [4], [18] To better simulate in-service conditions, G91 base metal plates were subject to PWHT for 2 hours at 760C. Table 4 shows the composition for the G91 plate used in this study.

23

Table 4: Compositions of G91 base metal and filler metals in wt%. Solid filler was used for CMT and FC Filler was used for FCAW as described below.

Wt C Cr Mo V Ni Mn N Nb B % Plate 0.08 8.40 0.92 0.23 0.09 0.45 0.041 0.075 10-4 Solid 0.106 8.92 0.99 0.194 0.47 0.75 0.044 0.063 3x10-4 Filler FC 1.4x10- 0.096 8.73 1.026 0.219 0.648 0.610 0.022 0.051 4 Filler d.1.2 Welding of G91 Two different welding processes were used to weld G91 specimens. The conventional flux-cored arc welding (FCAW) process is a constant power supply, continuously fed wire process. The wire is made up of a low-carbon steel shell and a flux-core which is used to alloy and shield the filler metal during welding. Figure 15a shows a schematic of the FCAW process. [74], [75] The second process is a discontinuously fed wire process known as cold metal transfer (CMT), in-which a solid filler wire is used. The CMT welding process uses wire- retraction to short the arc many times a second, which results in faster cooling-rates and lower heat-input than conventional wire fed arc welding processes. [72] Figure 15b shows snapshots from a high speed video from Pickin et al illustrating the CMT metal droplet process. [76] Filler wires are typically designed to match the creep-rupture strength of the BM. [18] G91 filler wire compositions used in the current study can also be found in Table 4. Welding was performed using a V-groove butt joint geometry, shown in Figure 16. CMT and FCAW welding was performed using a 20° included angle, on base metal plate with a thickness of 19.05 mm and a root opening of 25.4 mm. 13 mm thick G91 plate was used as weldment backing, as shown in Figure 16. Multipass welds were completed using 18-21 beads per weldment. Wired brushing was utilized between each pass to ensure slag removal and defect-free welds. Figure 17 shows a macrograph of an etched weldment, showing a typical bead sequence for each joint. [77] Pre-heat and interpass temperatures were maintained at 250°C and welds were shielded using an Ar-25%CO2 shielding gas. Pre-heat and interpass temperatures were monitored using thermocouples embedded in the backing plate. Pre-heat and interpass temperatures were maintained using a Harris pre-heating torch tip and oxy-acetylene fuel. [77] Following welding, all FCAW and CMT welds were post-weld heat treated (PWHT) for 2 hours at 760°C. To measure temperature profiles in the FGHAZ/ICHAZ of the CMT weldments, thermocouples were embedded in the base metal plate at various distances from the FZ to track heating rate, cooling rate and peak temperature as a function of distance from the FZ. Table 5 contains the welding parameters used to fabricate the CMT and FCAW welds.

24 Included in Table 5 is the calculated heat-input for both welding processes. The CMT welding heat-input is much lower than the FCAW heat input. Albert et al reported electron beam welding (EBW) and laser welding (LW) heat inputs for G91 weldments of 0.47 kJ/mm and 0.60 kJ/mm respectively which is similar to the heat-input produced by the CMT welding process. [6]

Table 5: Welding parameters used for CMT and FCAW welding processes.

Wire Feed Speed Travel Speed Travel Voltage Current Heat-input Process (mm/s) (mm/s) Angle (V) (A) (kJ/mm) CMT 35.84 1.44 10° push 13.4 168 0.67 FCAW 71.68 1.79 10° drag 25.0 238 1.40 d.1.3 G91 Simulated ICHAZ Specimen Preparation Simulated ICHAZ specimens were created using the thermal profiles shown in Figure 18, which emulate the thermal profile experienced in the ICHAZ during a single welding pass. Samples were produced using a GleebleTM 3500 thermo-mechanical simulator at OSU. The temperature profiles were chosen based on a combination of experiment and literature review. For the non-conventional CMT weldments, thermal profiles were measured during welding by embedding thermocouples into the weldment backing plate at various distances from the FZ to track heating rate, cooling rate and peak temperature as a function of distance from the FZ. For the CMT specimens, a heating rate of 800C/s, a cooling rate of 30C/s and a peak temperature of 900C was measured and used for the simulated ICHAZ specimens. For the FCAW specimens, values were extracted from existing literature on conventional high heat-input arc welding of G91 alloys. [78], [79] For the FCAW simulated ICHAZ specimens, a heating rate of 100C/s was utilized with a cooling rate of 10C/s and a 900C peak temperature. 10mm  x 63.5mm cylindrical samples were electro-discharge machined (EDM) for the Gleeble heat treatments. The EDM layer was ground off prior to processing to ensure good thermal conductivity between the Gleeble grips and the specimen surface. Figure 18 shows the thermal profile for each weldment. Temperature was measured and maintained using a type-K thermocouple welded to the sample surface at the center-length of the Gleeble specimens. After thermal processing in the Gleeble, specimens were machined for metallography or creep- fatigue testing. Figure 19 illustrates the experimental setup used to for the Gleeble ICHAZ simulations. d.2 G91 Creep-Fatigue Testing G91 base metal specimens for CF testing were extracted from the mid-section of the 19 mm thick plate with the loading direction parallel to the longitudinal rolling direction of the plate material . G91 weld metal specimens for CF testing were extracted with the loading direction perpendicular to the welding direction, including base metal, the entire HAZ and weld

25 metal in the gauge section of the CF specimens. Specimens for creep-fatigue testing were machined to the sample geometry found below in Figure 20, which conforms to ASTM E8 and ASTM E2714.[80], [81] The creep-fatigue testing apparatus was developed by retro-fitting a standard ATS-2400 series creep frame with a pneumatic actuator in the outer load-train. Using an electro- pneumatic regulator controlled by a PID controller written using LABVIEW system control software, the pneumatic actuator was filled and exhausted according to the desired loading waveform. Figure 21 shows a schematic of the cost-effective, load-controlled CF testing apparatus. Load, dwell time at maximum and minimum loads and loading rates were readily controlled using this testing method. Specimen displacement was measured using linear variable displacement transducers (LVDTs). All tests were conducted with a positive R-ratio, ranging from R=0.05 – R=0.14. Testing was conducted at 600C  2C and 650C  2C using a resistively heated clamshell style furnace and type-K thermocouples to monitor specimen temperature. Figure 22 shows a schematic stress-strain hysteresis and schematic loading waveform for the creep-fatigue tests conducted in the present study. d.5 Metallography Samples for optical and electron microscopy were prepared using standard grinding and polishing techniques. Samples were sectioned using a Struers Acutom-50 high speed saw. Deformed specimens were sectioned so the SEM viewing direction was perpendicular to the loading direction while the STEM viewing direction was parallel to the loading direction. Specimens were ground using 600, 800, and 1200 grit silicon carbide paper, followed by fine polishing using 6m, 3m, and 1m polycrystalline diamond paste on low-nap polishing cloths. A final polish was performed using 0.05m colloidal silica solution. Etching for optical microscopy on G91 specimens was performed using Vilella’s reagent, which consists of 5ml HCl, 1g picric acid and 100 mL ethanol. d.6 Characterization d.6.1 Optical Microscopy Several types of characterization were performed for the present study. Optical microscopy (OM) was completed on G91 weldments with Vilella’s reagent to calculate HAZ widths, locate specific regions in the HAZ and obtain macrographs of each of weldment. Optical microscopy was completed using a Klemex Digital Microscope. d.6.2 Scanning Electron Microscopy Scanning electron microscopy (SEM) was completed using a Thermo Fisher Apreo FEG

26 SEM at various beam conditions, depending on the desired information. The two primary imaging modes available in an SEM are formed with secondary electrons (SE) using an Everhart- Thornley detector or backscattered electrons (BSE) using a solid-state BSE detector. SE imaging collects low energy electrons and provides the user with topological information about the sample surface. In the current study, SE imaging mode was used for fractography and imaging etched specimens. BSE imaging mode collects high energy backscattered electrons to form images in which the main contrast mechanisms are the atomic number and crystallographic orientation of the interaction volume. BSE imaging mode was used in the current study to image the second phases present in CSEFS. General imaging in the current study was conducted at 5 kV accelerating voltage, 1.6 nA beam current, and 10 mm working distance. Using low accelerating voltage allowed the user to resolve the fine scale 100-200nm second phases present in CSEFS. Figure 23 shows schematic setups for SE and BSE imaging in an SEM. [82] d.6.3 Electron Backscattered diffraction Electron backscatter diffraction (EBSD) is an orientation imaging technique that enables determination of crystallographic orientation in bulk specimens. This involves collecting Kikuchi patterns formed by backscattered electrons to determine the crystallographic orientation of the interaction volume. Figure 24a below shows a schematic sample setup for EBSD, from Orsborn. [83] EBSD is used in the present study to measure high-angle (HAGB) and low-angle grain boundary (LAGB) content, and grain size in CSEFS before and after deformation. For EBSD, SEM operating conditions were set up at 20 kV, 3.2 nA beam current and 20mm working distance. For the examination of low-angle boundaries, a 3x3 3 pass Kuwahara filter was applied to the EBSD data. This is a common filter applied to EBSD data for the examination of substructure. [84] Transmission Kikuchi diffraction (TKD) is an orientation imaging technique performed using conventional EBSD hardware. To perform TKD, electron transparent specimens are prepared and positioned in the SEM such that forward scattered electrons from the incident beam are collected by the EBSD detector. Figure 24b shows the specimen and detector geometry utilized for TKD, from Orsborn. [83] This orientation imaging technique offers spatial resolution advantages when compared to conventional EBSD because the orientation information only comes from the final 5-10 nm of the thin foil. [85] Drawbacks of TKD compared to conventional EBSD include sampling limitations as well as extra time and costs associated with thin foil preparation. For TKD, high accelerating voltage is desired to get adequate signal through the thin foil. For the experiments in this study, TKD was performed using 30kV accelerating voltage and 6.4 nA beam current at 5 mm working distance. The sample was tilted to -20 with respect to the beam. To observe subgrains, TKD data was filtered using grain dilation and Kuwahara filtering techniques.

27 d.6.5 FIB Lamella Preparation To allow site specific sample extraction as well as to reduce the beam drift effects associated with conducting electron microscopy on ferromagnetic materials, specimens for STEM were prepared using a Thermofisher Helios NanoLab 600 Dual Beam Focused Ion Beam (FIB/SEM). Specimens were trenched and extracted from bulk specimens using 30kV ion beam accelerating voltage. Initial thinning was conducted at 30kV followed by final thinning at 5 kV. A final 900 eV argon mill was conducted on some FIB foils to remove surface damage associated with ion milling. d.6.6 Scanning Transmission Electron Microscopy Scanning transmission electron microscopy (STEM) was used in the present study to characterize precipitate type, distribution, and morphology for G91 specimens. In addition, STEM diffraction contrast imaging (STEM-DCI) was used to measure and quantify defect content such as dislocation density and subgrain size as a function of CF loading parameters, CF strain, and welding process. STEM offers many practical advantages over conventional TEM (CTEM). For example, STEM allows for the imaging of thicker samples, mutes bend contours, and allows the user to easily activate multiple diffraction conditions simultaneously. [86] Three main imaging modes were used to collect the STEM data presented in the present study: STEM brightfield imaging (STEM-BF), annular darkfield imaging (STEM-ADF) and high angular annular DF imaging (HAADF). STEM-BF and STEM-ADF are both useful for STEM-DCI. Intermediate convergence angles and larger camera lengths in STEM-BF and STEM-ADF mode maximize diffraction contrast and make it possible to image defects present in CSEFS, namely low-angle grain boundaries and dislocations. It is also possible to obtain images dominated by Z-contrast in STEM by using a HAADF detector. HAADF images are collected using large convergence angles and small camera lengths. Figure 25 illustrates the detector setup and schematic convergence angles used to conduct each type of STEM imaging. [86], [87] It is important to also consider camera length when choosing imaging conditions. Low camera lengths will cause highly diffracted electrons to be collected by the HAADF detector, optimizing Z-contrast. Higher camera lengths will cause electrons diffracted at lower angles to be collected by the ADF detector, maximizing diffraction contrast. STEM-ADF and STEM-BF imaging modes were used to image the different type of defects examined in the present study. Defect imaging was performed using a Thermofisher Tecnai F20 S/TEM operating at 200 keV accelerating voltage spot size 6. STEM-HAADF imaging was used in the present study to observe and quantify precipitate morphologies and distributions. STEM-HAADF imaging was performed on a Thermofisher Titan3 G2 60-300 S/TEM. In many cases it is necessary to determine the thickness of the volume being examined in the STEM. This was achieved by using a monochromated beam with a high current density to punch a hole through the thickness of the foil. After the cylindrical hole is formed the specimen is tilted to a large angle and the hole is viewed in projection. With a known tilt angle,

28 trigonometry is used to calculate foil thickness. Figure 26 shows a schematic of the process for thickness determination as well as images captured at each tilt angle. This method of thickness determination was used for all dislocation density measurements as well as any thickness measurements calculated from STEM images. d.6.7 Energy Dispersive Spectroscopy Energy dispersive spectroscopy (EDS) detects characteristic x-rays emitted from atoms during their interaction with the incident electron beam. EDS was used to qualitatively observe solute concentrations and identify second phases (carbides) present in the various CSEFS examined in the present study. Due to the low florescence- of light elements it was difficult to detect light elements such as carbon and nitrogen in the carbides present in CSEFS. However, heavier solute elements such as Cr and V were readily detected. EDS was conducted on a Thermofisher Titan3 G2 60-300 S/TEM equipped with a chemiSTEM quad silicon drift detector system. The microscope was operated at 300kV and spot size 6 for EDS collection. To maximize EDS counts, current density was increased by condensing the monochromator above what would be used during normal imaging conditions, typically around 0.3-0.4 nA.

e. G91 Base Metal Characterization and Creep-Fatigue Studies e.1 Introduction and Objectives for Base Metal CF Studies The ferritic-martensitic steel alloy 9Cr-1MoV (G91) has a high thermal conductivity, low coefficient of thermal expansion, and excellent creep strength and oxidation resistance, making this material advantageous for elevated temperature use. [3]–[5], [37] Typical components made from G91 include heat exchangers, pressure vessels and piping.[3] G91 is also a candidate material for the next generation of sodium-cooled fast reactors (SFRs). In-core applications include cladding wrappers and ducts. Current temperature ranges for G91 are approximately 540C - 600C.[1], [3] While in service, G91 components are exposed to elevated temperatures and cyclic stresses, leading to creep-fatigue (CF) damage accumulation. While there are a multitude of creep and low-cycle fatigue studies available for G91, there is a knowledge gap concerning coupled, creep-dominated creep-fatigue data.[16], [32], [51] The majority of creep-fatigue studies available are strain-controlled, also known as stress-relaxation fatigue (RF) tests, which may not accurately reflect in-service conditions, and also make it difficult to study the effect of anelasticity on creep-strain accumulation. [37], [40], [42], [48] One reason the majority of available creep-fatigue data is collected in a strain-controlled, RF testing mode is the fact that most data is collected on servo-hydraulic test machines, which lend themselves to displacement-controlled testing modes. For the current study, a custom- built, load-controlled creep-fatigue testing apparatus was constructed in an effort to conduct creep-dominated creep-fatigue tests. Performing load-controlled testing also allows for the study of anelastic backflow of strain at minimum load dwell periods, which has been known to occur in alloys that develop substructure. [46], [88], [89] Many of the available studies focus on

29 life-prediction via the collection of mechanical data, without regard to microstructural progression of these alloys during CF or RF deformation. [15], [40] The current study attempts to bridge this knowledge gap by conducting interrupted CF tests to study microstructural evolution as a function of CF strain accumulation. G91 components are employed in the normalized and tempered condition, consisting of an austenitization treatment at 1040C, followed by air cooling to produce a fully martensitic microstructure. The material is then heat-treated (typically for 1 hour) at 760C to temper the martensite and precipitation harden the matrix. [4] The resulting microstructure consists of prior austenite grains (PAGs) containing high-angle block and block boundaries as well as low- angle lath boundaries. The martensitic transformation is accompanied by a large strain energy 14 -2 which produces significant dislocation density (dis >10 m ) that persists through tempering. The large dislocation density along with low-angle lath boundaries create an extensive subgrain network in CSEFS. During tempering, second phases precipitate at high and low-angle boundaries. The dominate second phase in 9-12% Cr steels is M23C6. Second phases are thought to play a critical role in the monotonic creep strength of these alloys by pinning the substructure and preventing coarsening during creep. [29], [34] In addition, during austenitization and subsequent tempering, V-rich MX particles nucleate within subgrain interiors, inhibiting mobile dislocation motion during elevated temperature deformation. Figures 27 and 28 show the microstructure of G91 in the as-tempered condition. Figure 27 shows the characteristic microstructure of G91 in the normalized and tempered condition at intermediate magnifications. A mixture of high-angle (prior austenite grain boundaries, block boundaries, and block boundaries) and low-angle (subgrain boundaries made up of martensite laths and dislocation arrays) are decorated by M23C6 carbides. Figure 29 shows the microstructure of G91 at higher magnification using STEM-DCI. Fine subgrains are decorated with M23C6 carbides which are responsible for pinning low-angle boundaries. High dislocation density and MX precipitates, which are responsible for pinning the mobile dislocations exist within the subgrain interior. Figure 29 also shows an example of the MX carbonitrides present in V-added CSEFS like G91. MX particles exist within the subgrain interior with a spherical morphology and inhibit mobile dislocation motion during CF deformation. Elevated temperature strength in G91 and other CSEFS is derived from multiple strengthening mechanisms. The primary strengthening mechanism is sometimes referred to as “carbide-stabilized substructure hardening”, in which substructure coarsening is prevented by precipitation of pinning particles at subgrain boundaries. [29], [34] In addition, MX particles inhibit recovery of mobile dislocations, and Mo, Cr and Mn contribute to solid solution strengthening. [3], [33] While the monotonic creep properties of G91 have been studied extensively, a knowledge gap pertaining to creep-dominated, creep-fatigue deformation exists in CSEFS. In recent reporting by power generation authorities, creep-fatigue deformation was recognized as a critical design issue for high-temperature structural design for power applications, with

30 emphasis on the lack of available creep-fatigue data and mechanistic knowledge concerning CF damage in G91 and other CSEFS components. [16] Figure 30 shows a comparison of load- controlled creep-fatigue data collected at multiple temperatures and loads compared to monotonic time-to-rupture values from data collected by Cerri et al. and Tabuchi et al. [32], [51] From Figure 30 it is clear that the time-to-rupture for CF specimens is significantly shorter at all test conditions. This proof of concept illustrates the deleterious effects associated with the addition of a fatigue component to elevated temperature structural components made from CSEFS. The current study attempts to examine the effects of loading parameters and temperature on mechanical response for load-controlled, creep-dominated CF of G91. In addition, a comprehensive assessment of microstructural evolution during creep-fatigue deformation is examined with special emphasis on microstructural features that have been found to be critical to elevated temperature strength of G91. This section examines the initial microstructure of G91 specimens after normalization and tempering. The creep-fatigue properties of G91 are examined, including mechanical response under various loading conditions at 600C and 650C. Microstructural evolution as a function of creep-fatigue damage is also studied. A load-controlled testing apparatus was designed and built to conduct load-controlled, creep dominated creep-fatigue testing. Quantitative microstructural studies found substructure coarsening and a drastic decrease in dislocation density after creep-fatigue deformation. Significant amounts of anelastic backflow were observed at minimum load during every creep-fatigue test conducted. The effects of loading parameters on anelastic backflow were also studied. Fracture occurred in a ductile manner for every specimen examined. A high density of creep-voids was found in fractured specimens at grain boundaries and precipitates. Comparisons of creep-fatigue data collected in the present study to previous monotonic creep studies found that the same operative mechanisms control deformation during monotonic creep and creep-fatigue. However, subgrain coarsening occurs much more rapidly under CF conditions than under monotonic creep conditions, resulting in higher strain-rates and shorter rupture lives under creep-fatigue conditions, compared to monotonic creep. e.2 Experimental Procedures Plates of G91 with a 19 mm thickness were received from American Alloy Steel in the normalized and tempered state. The plates were austenitized at 1040C, followed by air cooling and a subsequent 1-hour tempering treatment at 760C. Plates were heat treated again at 760C for 2 hours to simulate in-service conditions which would typically include a PWHT. CF specimens were machined using the specimen geometry found in Figure 20. Specimens were tested on the custom-built CF testing apparatus described in section d.2. Table 6 contains the testing parameters for all base metal tests conducted in the current study.

31 Table 6: Summary of base metal CF tests conducted in the current study.

Samples for microscopy were prepared using conventional metallography techniques. SEM imaging was performed at 5kV and 1.6nA. EBSD was performed at 20kV accelerating voltage and 3.2 nA beam current. TKD was performed at 30kV and 6.4nA beam current. Samples for microscopy were prepared using FIB for site specific sample extraction and also to reduce the ferromagnetic effects associated with electron microscopy of ferromagnetic materials. STEM imaging was performed at 200 kV and spot size 6. M23C6 precipitate distributions were obtained by segmenting SEM-BSE images using MIPAR image processing software. e.3 Results and Discussion e.3.1 Initial Microstructure The microstructure in the as-received material, examples of which can be found in Figures 27 and 28, was consistent with previous studies. [30], [57] Marked features include high dislocation density (2.92 x 1014 m-2) and a fine subgrain network. Initial subgrain measurements suggest a mean equivalent diameter of 547 nm  45 nm. Dislocation density measurements were calculated using low-order zone-axis STEM images. Zone-axis diffraction- contrast imaging methods make it possible to activate multiple diffraction conditions simultaneously, providing a good overall view of dislocation content within each sample. Equation (5) was used to measure the dislocation density in the as-received and deformed specimens. This stereological method for determination of dislocation density is widely accepted and is described in detail in the included references. [30], [90]

1 푁퐿 푁푇 𝜌푑푖푠 = ( + ) (5) 푡 퐿퐿 퐿푇 -2 where ρdis (m ) is the dislocation density, t (m) is specimen thickness, NL and NT are the

32 longitudinal and transverse dislocation intersections, respectively, and LL (m), and LT (m) are the longitudinal and transverse line lengths, respectively. Figure 31 shows a STEM-ADF image from a ruptured G91-CF specimen tested at 650C and 100 MPa max load with a 15 minute dwell time. The image in this figure was taken on a [111] zone axis and has been overlaid with a grid to illustrate the stereological method used to calculate dislocation density. Some of the dislocation intersections are marked as well. e.3.2 Microstructural evolution with CF strain A representative load-controlled CF curve from a G91 specimen is shown in Figure 32. The first feature to note is the CF curve’s resemblance to a traditional monotonic creep curve. The CF curve in Figure 32 shows a clear primary, secondary and tertiary regime with strong primary and tertiary transients, all features that are commonly used to describe monotonic creep deformation. The 1% strain accumulation during the primary creep-fatigue transient is indicative of hardening, which is commonly found in class II alloys under monotonic creep deformation conditions. Traditionally, class II alloys exhibit large primary transients due to the formation of subgrain networks which stabilize the microstructure and result in a reduction in monotonic creep strain-rate. [91] This is not the case for G91 which starts out with a large dislocation density and refined substructure. [33], [34] High dislocation density and refined substructure in the tempered material are the direct result of the martensitic transformation that occurs in high Cr CSEFS. [3], [30] During tempering, laths made up of low-angle boundaries become equiaxed, and there is rearrangement of free dislocations into low energy substructures. In Figure 32, the overlaid solid markers denote interrupted CF tests conducted with identical testing parameters. The open circle in this figure denotes a sample extracted from the failed specimen at a distance from the fracture surface that would emulate the strain accumulated at the desired cycle. To the author’s knowledge this is the first study observing microstructural evolution in a G91 during load-controlled creep-fatigue deformation. Subgrain size as a function of CF cycle with an overlay of strain rate vs cycle number for the failed specimen is shown in Figure 33. There is slight coarsening during the primary transient, however, as the CF strain rate reaches a minimum and begins to increase, subgrains coarsen rapidly until failure. Subgrain coarsening is due to the migration of low-angle boundaries upon the application of stress at elevated temperature. Substructure coarsening is a well-documented phenomenon during monotonic creep of CSEFS. [33], [63], [92], [93] The only way for appreciable substructure refinement to occur is by the formation of new subgrain boundaries. The original substructure present after tempering is primarily due to the martensitic transformation, however, during elevated temperature deformation, new subgrains form from the rearrangement of dislocations into low energy structures. Typical class II alloy primary creep is the result of hardening by the rearrangement of dislocations into low energy structures via the formation of subgrains. [91] CSEFS are unique in that they begin with very

33 high dislocation densities and refined subgrain networks. The result of the starting microstructure is a competition between substructure refinement from the rearrangement of dislocations and substructure coarsening due to bowing and migration of subgrain boundaries. Simultaneous subgrain coarsening and rearrangement of dislocations into low energy structures results in the initial hardening and also delays the coarsening observed in specimens with more CF strain. These dynamic changes in subgrain size are illustrated by the subgrain size evolution shown in Figure 33. Examples of STEM diffraction contrast images (STEM-DCI) used calculate subgrain size are shown in Figure 34. Subgrain mean equivalent diameters were determined using the line intercept method which is outlined in ASTM E112. [94] STEM-DCI images of representative examples of subgrain size as a function of CF cycle number are shown in Figure 34. During the primary and secondary regimes of the CF test, subgrain boundaries are decorated with Cr-rich M23C6 carbides. However, after failure the M23C6 carbides are no longer acting as effective obstacles to subgrain motion. Many of the M23C6 particles are located in the subgrain interior after CF rupture. CSEFS exhibit a large dislocation density after tempering, a remnant of the martensitic transformation upon cooling from austenitization. [33] Along with subgrain size evolution, understanding the variation in dislocation density as a function of CF strain is critical to the underlying mechanisms governing creep-fatigue deformation in CSEFS. Figure 35 shows dislocation density (not associated with subgrain boundaries—thus, the “mobile” dislocation density) as a function of cycle number for the interrupted CF tests. After an initial rapid decrease in dislocation density, subgrain coarsening and a more gradual decrease in dislocation density occur simultaneously until failure. The large dislocation density in the undeformed material most likely causes the significant primary transients found in most CSEFS under CF and monotonic creep deformation. [95] Dislocation density was measured using zone-axis diffraction contrast imaging. Figure 36 shows examples of low-order zone-axis images used to calculate dislocation density. Examples of boundary dislocations are indicated by arrows in the zone-axis STEM-DCI images from the undeformed base metal and the specimen extracted at cycle 26. For the current analysis, boundary dislocations are not considered when measuring the “mobile” dislocation density. While the diffraction-contrast attainable using STEM-DCI provides an excellent way to observe and measure subgrain evolution during aging and elevated temperature deformation, the drawback of this experimental method exists due to the limited sampling area possible. To complement the STEM based subgrain measurements, EBSD was conducted on the interrupted specimens, allowing for the observation of substructure over larger areas. The drawback of conventional EBSD lies in its angular resolution limits, which make it difficult to resolve subgrain boundaries with very small boundary misorientations. To resolve the substructure present in the EBSD scans, data was run through an edge preserving Kuwahara filter using 3 passes of a 3x3 grid. This filtering technique has previously been used to resolve subgrain boundaries in deformed materials. [84], [96] Figure 37 shows EBSD grain boundary maps overlaid on image quality (IQ) maps to illustrate substructure evolution as a function of CF strain accumulation, the low-angle grain boundary (LAGB) content in the grain boundary maps aligns well with the

34 coarsening characteristics observed from the STEM-DCI measurements. For the present study, LAGB content includes any boundary with a misorientation less than 15. Figure 38 illustrates the total LAGB length as a function of CF cycle number. The LAGB content is plotted as 1º-5º misorientation and 5º-15º misorientation. From this figure it is clear that the majority of coarsening occurs for the lowest angle boundaries and the 5º-15º boundary content remains relatively constant. It has been shown for monotonic creep that substructure hardening is the most significant strengthening mechanism for CSEFS. [29], [34] The evolution of subgrain size has been studied extensively for CSEFS under monotonic creep conditions.[33], [63], [92], [93] Figure 39 shows subgrain size measurements taken from the present interrupted creep-fatigue tests together with subgrain size measurements from monotonic creep experiments conducted by Maruyama et al. and Aghajani et al. [33], [92] The subgrain coarsening data in Figure 39 from monotonic creep experiments was conducted on various 9-12 wt% CSEFS including G91 and X20 at 650C and 550C. From this figure it is clear that subgrain coarsening occurs significantly faster for the creep-fatigue specimens than for the monotonic creep specimens. The best comparison in Figure 39 is made when comparing the present CF subgrain coarsening measurements to the G91 monotonic creep data conducted at 650C and 98 MPa maximum load by Maruyama et al. [33] This data shows that even while at higher maximum load, the monotonic G91 specimens show significantly decreased coarsening rates when compared to the CF subgrain coarsening measurements. While final subgrain size is fairly consistent regardless of the method of deformation, subgrains coarsen more quickly under CF deformation conditions. Cerri et al. and Eggeler et al. showed that particle coarsening is heavily dependent on monotonic creep conditions and coarsening is accelerated during creep compared to stress-free aging at equivalent temperatures. [32], [34] Subgrain coarsening becomes more severe as particles coarsen. M23C6 precipitates are responsible for pinning the subgrains, and as they coarsen, their potential for pinning decreases. Examples of low-angle boundaries being pinned by M23C6 precipitates are shown in Figure 40. Coarsening of boundary precipitates results in a decrease in particle density along low-angle grain boundaries and a decrease in pinning pressure per particle. The Zener pinning force varies inversely with particle diameter, as shown in Equation (6). [97] 3푓훾 퐹 = (6) 푍푒푛푒푟 2푟 where FZener (N) is the Zener pinning force, 푓 is the boundary particle volume fraction, 훾 is the interfacial energy (J/m2), and r (m) is the particle radius. Figure 41 illustrates the cumulative carbide size probability for G91 CF specimens as a function of CF cycles. M23C6 mean equivalent diameters were generated by segmenting SEM- BSE images, like the one shown in Figure 27a. Figure 41 shows that the most substantial coarsening occurs at the later stages of CF strain accumulation.

35 e.3.3 Effect of Loading Parameters on Minimum Strain-rate e.3.3.1 Effect of Load on Minimum Strain-rate Extensive studies have been conducted to determine the creep mechanisms governing monotonic creep deformation in CSEFS. [32], [33], [98] One classic way to distinguish between different operative creep mechanisms is by developing Norton-power law parameters.[99] The phenomenological power-law relationship allows experimenters to distinguish between different creep mechanisms by a change in the stress exponent in Equation (7). [99]

A 휎 푛 휀̇ = ( ) (7) 푚푖푛 kT 퐺 where 휀푚푖푛̇ (1/s) is the minimum strain rate, A and n are material constants, k (J/K) is Boltzmann’s constant, T (K) is temperature, 𝜎 (MPa) is stress, and G (Pa) is bulk modulus. Due to the creep-dominated nature of the testing conducted in this study, it is logical to evaluate the data collected using a power-law relationship. Figure 42 shows the power law exponents derived from creep-fatigue data collected at different loads, with a dwell time of 15 minutes. A stress exponent of n=6.9 was determined for G91 at 650C under matching CF loading conditions. At 600C, a stress exponent of n=8.9 was determined. While the stress exponent may be dependent on maximum and minimum load dwell time, the stress exponents for a 900s max load dwell time match those reported for monotonic creep well, as illustrated by the monotonic creep data extracted from the studies of Cerri et al and Cadek et al. [32], [100], [101] The data from Cerri et al and Cadek et al suggest n=8 at 650C and n=12 at 600C. The general agreement of the stress exponents reported in existing monotonic creep literature and those produced from creep-fatigue data collected in the current study, indicates that the same mechanisms are governing CF deformation and monotonic creep deformation. However, from the time-to-rupture comparisons illustrated in Figure 30 as well as the subgrain coarsening comparison shown in Figure 39, it is clear that CF is a much more severe mode of deformation than monotonic creep, causing higher minimum strain rates, accelerated subgrain coarsening and shorter rupture times than monotonic creep tests with the same maximum applied stress. e.3.3.2 Anelastic Backflow Materials that develop substructure during elevated temperature deformation, commonly referred to as class II alloys, have been known to exhibit anelastic backward flow of strain after abrupt stress reductions during creep deformation. [46], [47] Sawada et al found this to also be the case for CSEFS, however, CSEFS begin with a refined subgrain network, as opposed to the formation of substructure after the onset of creep, which is the case for conventional class II pure metals and alloys. Sawada et al’s experiment consisted of abrupt stress reductions on alloy P92 at elevated temperatures, tracking displacement as a function of loading and time. [48] For a load-controlled creep-fatigue experiment, this process is repeated regularly throughout the course of the test, as exhibited by the waveform in Figure 22. As a result, during the unloading and minimum load portion of each cycle, some anelastic backflow

36 of strain occurs. The anelastic backflow in the present CF testing is demonstrated by Figure 43, which contains three stress-strain hysteresis loops extracted from a creep- fatigue test during the primary, secondary, and tertiary regimes of the experiment (strain in each loop was smoothed with a moving average filter). The amount of “forward” creep-strain accumulated at maximum load varies significantly as a function of cycle number/total strain accumulation. The change in “forward” creep strain per cycle is consistent with the variation in the overall minimum creep rate shown in Figure 33. However, the amount of anelastic backflow during unloading and at minimum stress is relatively constant as a function of cycle number. A series of CF tests were conducted to study the effect of CF loading parameters on anelastic strain. Figure 44 shows the creep strain and anelastic strain per cycle from the creep- fatigue test shown in Figure 32. While the creep strain per-cycle experiences significant primary and tertiary transients, the amount of anelastic backflow marginally decreases as a function of cycle number. CF loading parameters have a significant effect on the amount of anelastic backflow per cycle. Figures 45 and 46 show the average anelastic strain per cycle during the secondary regime as a function CF loading parameters. Figure 45 illustrates the effect of dwell time at maximum load and dwell time at minimum load on anelastic strain. As might be expected, as the time spent at minimum load (10 MPa for every test in this study) increases, the average anelastic strain accumulated per cycle increases as well. In addition, as the dwell time at maximum load increases, anelastic strain increases. In addition to studying the effects of dwell time on anelastic strain, the effect of creep load on anelastic strain was also investigated. Figure 46 illustrates the effect of creep load vs anelastic strain for specimens tested with the same dwell time parameters. From this figure it is clear that as the load is increased, the amount of anelastic backflow is also increased. Figures 45 and 46 suggest that as creep strain per cycle increases, the amount of anelastic backflow per cycle also increases. Also shown in Figure 46 is rupture time as a function of creep load and anelastic backflow. As expected, as creep load is increased, rupture life is decreased. Understanding the mechanisms governing anelasticity after stress-reductions in CSEFS is critical to understanding load-controlled CF deformation. Anelastic backflow must come from the dense substructure present in all CSEFS after tempering. Gibeling and Nix showed that a variety of FCC class II alloys exhibited anelastic, time-dependent recovery during abrupt stress reductions after substructure had formed in the alloys during creep. [47] In addition, Gibeling and Nix concluded that anelastic backflow of strain was due to internal back-stress from the microstructure. [46], [47] Sawada et al found that Cr-Mo steels did not exhibit anelastic strain upon unloading unless the alloy had undergone a martensitic transformation and exhibited a lath structure, which consists primarily of low-angle boundaries. They also determined that backflow was present regardless of dislocation density. [48] There is sufficient evidence to support the role of subgrain structure in the anelastic backflow that occurs at minimum load during CF testing. Sawada et al attributed this anelastic strain to bowing of subgrain boundaries during creep, which led to anelastic relaxation after stress-reductions in Cr-Mo steel specimens. They also showed that steels which exhibit less

37 anelastic backflow also exhibit improved monotonic creep-strength, associating the difficulty of boundary movement via carbide stabilization with reduced anelasticity and increased creep strength. [48] Repeated forward bowing and anelastic backward bowing during cyclic loading may accelerate subgrain coarsening and rupture, compared to monotonic creep deformation, in which the flow of strain is quasi-static and entirely in one direction. In addition, it has been noted that the microstructure of CSEFS is extremely inhomogeneous, especially after the onset of creep damage. [30] Such inhomogeneities consist of subgrains with large dislocation densities adjacent to subgrains with similar orientation but vastly less dislocation content. The current study has shown that the heterogeneous microstructures after creep are also present after CF deformation. An example of the heterogeneous microstructure is shown in Figure 47. Figure 47a shows a zone-axis diffraction contrast image of two subgrains with similar orientation but extremely different dislocation densities. Figure 47b and c show transmission Kikuchi diffraction (TKD) data illustrating the similar orientation of the subgrains in an IPF map but varying lattice strain in the grain reference orientation deviation (GROD) map of Figure 47c. These heterogeneous microstructures present in ferritic-martensitic steels may also contribute to the anelastic strains exhibited during load-controlled creep-fatigue and abrupt stress reductions during monotonic creep. One study by Gibeling and Nix was based on the hypothesis that a heterogeneous microstructure was produced, consisting of hard subgrain walls and soft subgrain interiors, which lead to the generation of internal stresses. This concept could also be applied to hard and soft subgrains (or subgrains with high dislocation density and subgrains with low dislocation density). The anelasticity may be a cumulative effect of anelastic bowing of subgrain walls as well as load shedding from soft subgrains to strain-hardened subgrains. Gibeling and Nix showed that anelastic backflow from subgrain bowing could produce negative strain with similar magnitude to elastic strains accumulated during loading. [46], [47] Sawada et al showed that CSEFS which did not contain a refined subgrain network did not experience anelastic backflow upon abrupt stress reductions, regardless of mobile dislocation density. [48] Therefore, it is reasonable to assume that the majority of anelastic backflow during CF in CSEFS is due to the reverse bowing of subgrain boundaries during unloading and minimum load. The most significant difference between monotonic creep deformation and creep-fatigue deformation is the repeated forward and backward motion of subgrains during creep-fatigue, compared to the quasi-static forward subgrain migration during monotonic creep. It is also reasonable to expect anelastic backflow to be closely related to the accelerated subgrain coarsening exhibited under CF deformation, compared to the relatively slow subgrain coarsening rates under monotonic creep deformation at comparable loads. STEM-DCI in combination with SuperX EDS performed on deformed CF specimens show evidence of solute enrichment at subgrain boundaries in G91 specimens. EDS maps from an edge-on subgrain boundary in the specimen interrupted at cycle 26 (shown in Figure 32) are shown in Figure 48. The EDS maps show chromium and molybdenum enrichment and iron deficiency at the subgrain boundary. The holes shown on each side of the subgrain boundary are nanoholes drilled with a monochromated beam to serve as fiducial markers for EDS map

38 drift correction. Subgrain boundaries are made up of dislocation arrays or networks (an example is shown in Figure 40). Therefore, it is reasonable to assume the solute observed in the EDS maps in Figure 48 are solute atmospheres around the individual dislocations which make up the subgrain boundary. As the subgrain migrates under an applied stress during monotonic creep deformation, the subgrain will experience an effective stress which includes a positive contribution from the applied stress, and negative contributions from solute drag as well as obstacles such as M23C6 particles. However, in the cyclic case of CF deformation, the subgrain boundary may break free from the solute during the rapid anelastic motion which occurs during unloading and while at minimum load. Following the reverse boundary motion, the solute begins to diffuse away. Upon reloading, the subgrain boundary can migrate by thermally activated climb and glide without the solute drag effect for some period of time, until the solute is recollected. Initial measurements using a 푥 = √퐷푡 relationship with diffusion coefficients calculated using Thermocalc indicate a diffusion distance of 2.9 nm for Cr and 3.5nm for Mo during a 10s minimum load condition. This period of accelerated subgrain boundary migration may lead to accelerated subgrain coarsening and CF rupture. This hypothesis is illustrated in Figure 49. Monotonic creep and the later portions of the CF dwell period are represented by Figure 49B, where the individual dislocations in a simple tilt boundary each contain a solute atmosphere. Following unloading, the subgrain breaks free of the solute atmosphere, as shown in Figure 49C. After the boundary dislocations break free, the solute may diffuse away and upon reloading, the subgrain boundary migrates without the solute drag effect (represented by Figure 49D), until the solute is recollected and solute drag resumes (represented by Figure 49B). More work is necessary to determine the legitimacy of this hypothesis. Principally, it is necessary to determine the magnitude of forward and reverse subgrain boundary motion during each cycle, which may require modeling efforts. Gibeling and Nix have produced analytical models for subgrain boundary motion in pure metals, under monotonic creep conditions. [46] In addition, Ma and Wang et al have modeled the effect of solute drag on high-angle grain boundary migration.[102] These previous modeling efforts may provide a framework for modeling the effect of solute drag on anelastic subgrain boundary motion and accelerated subgrain coarsening in G91 under CF deformation conditions. e.3.3.4 Effect of dwell Time on CF deformation The current study also attempts to examine the effect of dwell time on CF deformation. Figure 50 shows four CF tests conducted at 650C, 85 MPa maximum load and 10 MPa minimum load. The only varied loading parameter between the specimens is dwell time which was varied between 300s-5400s. After examining the effect of creep load on rupture life shown previously in Figure 46, one might be expected that rupture life should vary inversely with increasing dwell time, since the amount of creep-strain per cycle increases with increasing load and increasing dwell time. However, Figure 50 illustrates that the dwell time does not have a clear relationship with rupture life, as the specimen with the intermediate 3600s dwell time

39 failed after fewer cycles the longest dwell time of 5400s. To understand how maximum load dwell time affects rupture life it is important to observe the effect of dwell time at maximum load on anelastic backflow. Figures 51 and 52 show the anelastic backflow per cycle and the creep strain per cycle respectively. From these two figures it is clear that as the maximum load dwell time increases, anelastic backflow and creep strain per cycle increase as well. Subtracting the backward flowing anelastic strain from the forward flowing creep strain allows for the calculation of net forward flow of strain per cycle. The average anelastic strain, creep strain, and net forward flow of strain for the secondary regimes of the CF specimens with varied maximum load dwell times are shown in Figure 53, plotted as a function of dwell time. From this figure it is clear that anelastic backflow increases at a greater rate than creep strain (over this range of dwell times), as maximum load dwell time is increased. This potentially accounts for the counterintuitive relationship between rupture life and maximum load dwell time. An examination of average net strain per cycle as a function of cycles to failure for specimens with varied maximum load dwell time is shown in Figure 54. It shows that as net strain per cycle increases, cycles to failure decreases. e.3.4 Creep-Fatigue Fracture Characteristics Fractography performed on G91 CF specimens revealed coalescence of creep voids as the main fracture mode. Figure 55A shows a fracture surface exhibiting ductile fracture characteristics and void coalescence. All specimens failed in a ductile manner exhibiting significant necking and coalescence of creep voids. Figure 55B shows a polished fracture surface, showing coalescence of creep voids. Figure 55C and D show cross-sections of a failed CF specimen. Creep voids are present at grain boundaries and precipitates and grow with decreasing distance to the fracture surface. This type of failure is indicative of monotonic creep failure in CSEFS. [57], [64], [103] However, creep-fatigue specimens tested using strain- controlled RF methods exhibit failure characteristics that are typical of pure-fatigue experiments, including microcracking and fatigue striations. [37], [38], [40], [42] Fracture of G91 specimens under load-controlled CF conditions is more comparable to monotonic creep- rupture than failure under pure-fatigue conditions. e.4 Summary and Conclusions of Base Metal CF Studies The present work illustrates the deleterious effects of combined creep-fatigue, compared to monotonic creep. Microstructural evolution and the effect of loading parameters and temperature on mechanical response are examined. A s a result, the following conclusions are made:

40  After tempering, CSEFS contain a high dislocation density, refined subgrain size and fine distribution of particles along LAGBs and HAGBs, all of which continuously evolve during CF deformation.  Studying microstructural evolution reveals that subgrain coarsening is sluggish at the onset of CF damage but occurs rapidly in later stages of CF deformation. Cooperatively, dislocation density decreases rapidly at the onset of CF damage and decreases at a continuously slower rate as a function of CF cycling.  A significant amount of anelastic backflow of strain occurs during every load-controlled CF cycle. Anelastic backflow increases with time spent at minimum and maximum load, as well as increasing creep load. Anelasticity is thought to be caused primarily by subgrains bowing around pinning precipitates.  Fracture occurred in a ductile manner for all specimens tested, fracture characteristics include creep void formation and coalescence at the fracture surface. A high density of creep voids was found at grain boundaries and precipitates. Failure characteristics were more comparable to those reported in monotonic creep studies than fractography performed in pure-fatigue or RF studies found in existing literature.  At the testing temperatures of 600 and 650 C, the same mechanisms are operating during load-controlled CF and monotonic creep, producing similar stress exponents and microstructural response to loading. However, subgrain coarsening occurs more quickly under CF loading conditions, resulting in higher strain-rates and shorter rupture lives. It is postulated that the repeated backward and forward subgrain bowing that occurs during loading, unloading and at minimum load reduces substructure stability and accelerates subgrain coarsening, leading to increased strain-rates and shorter rupture lives when compared to monotonic creep experiments.

f. Creep-fatigue Deformation and Characterization of G91 Weldments f.1 Introduction and Objectives for G91 Weld Studies Microstructure of G91 in the normalized and tempered condition is shown in Figure 56. As discussed previously, elevated temperature strength in G91 is derived from a combination of the following strengthening mechanisms: a carbide-stabilized substructure strengthening effect by the M23C6 precipitates, a strain-hardening effect from high dislocation density and a precipitate strengthening effect from MX particles pinning mobile dislocations. [3], [29], [33], [34], [105] During welding, the base metal microstructure is markedly altered by the welding thermal cycles. An optical micrograph spanning the HAZ in a G91 weldment can be found in Figure 57. The coarse-grained heat-affected-zone (CGHAZ) reaches a temperature well above A3 during welding and all carbides present prior to welding are dissolved, resulting in large prior austenite grains (PAGs) after cooling. The FGHAZ also reaches a temperature above A3, however, carbides are only partially dissolved, resulting in a region of fine PAGs, where grain

41 size is limited via grain boundary pinning by carbides. The ICHAZ reaches a temperature between A1 and A3 resulting in the formation of fresh martensite, tempered martensite, and coarsened carbides after welding. The performance of G91 is severely restricted by frequent premature failures in the welded sections, specifically in the heat-affected-zone (HAZ) of the weldment. [5] Creep and CF damage is especially detrimental to G91 weldments, resulting in premature failure, typically occurring in the fine-grained or inter-critical heat affected zone (FGHAZ/ICHAZ). [16], [17], [104] Such failure in the outer region of the heat-affected zone is known as type IV failure. [18] As reviewed in Section c, many different mechanisms have been proposed to explain Type IV failure, including changes in precipitate distribution, size and morphology, ferrite formation in the FGHAZ during cooling, and stress triaxiality from the microstructural and corresponding mechanical property gradient created by the thermal gradient in the HAZ during welding. [6], [7], [10]–[13] Research efforts to prevent type IV failure and improve weldment strength have been promising. Previous studies have shown that by using low-heat-input welding processes such as electron beam (EBW) and laser welding (LW), the width of the HAZ can be reduced and the creep rupture life of G91 weldments increased. [6] Multiple mechanisms have been suggested to account for the increase in monotonic creep strength from low-heat-input welding processes, the two major theories are (1) a change in precipitation behavior in the HAZ and (2) a change in stress triaxiality in the heat affected zone leading to increased HAZ creep strength. It should be noted that the welding processes used to create low-heat-input welds in previous studies were EBW and LW, processes that would be impractical for implementation in a service environment. [6] However, a recently developed arc welding process known as cold metal transfer (CMT) welding uses a pulsed arc to deposit weld metal into the weld joint, resulting in significantly decreased heat-input compared to processes conventionally used to weld G91 components, such as flux-cored arc welding (FCAW). [72], [106] To the author’s knowledge, this is the first study to systematically compare the elevated temperature properties of G91 joined using the CMT welding process. The objectives of this study are to evaluate the microstructural and mechanical differences between G91 weldments created using a conventional high-heat input process and a non-conventional low-heat input process that could be field applicable. In addition, the mechanisms behind the significant differences in elevated temperature strength are investigated. Specimens from each welding process were tested using the custom-built creep-fatigue testing apparatus under multiple maximum loads at 650C. Ruptured weldments were examined using characterization techniques including scanning and transmission electron microscopy (SEM/STEM). S pecimens welding using the CMT welding process significantly outperformed FCAW weldments. Further characterization revealed that changes in precipitate size and distribution as well as differences in subgrain size and dislocation density between the two welding processes resulted in the differences in creep-fatigue strength.

42 f.2 Materials and Experimental Procedures G91 base metal was received in plate form, in the normalized and tempered state. Flux- cored filler wire was used for the FCAW weldments and solid filler wire was used for the CMT weldments. Filler wires are typically designed to match the creep-rupture strength of the BM. [18] Table 4 summarizes the G91 base metal and filler wire compositions used in this study. Welding was performed using the procedures described in section d. Table 5 contains the welding parameters used to fabricate the CMT and FCAW welds. Included in that table is the calculated heat-input for both welding processes. It is clear that the CMT welding heat-input is much lower than the FCAW heat-input. Albert et al reported electron beam welding (EBW) and laser welding (LW) heat-inputs for G91 weldments of 0.47 kJ/mm and 0.60 kJ/mm respectively. [6] Comparison of the CMT heat-input to literature values for EBW and LW show that the CMT arc welding process produces welds with similar heat-input. Samples for characterization were prepared using conventional metallographic techniques. Samples for microscopy were produced by polishing using 1200 grit SiC followed by 6 µm, 3 µm and 1 µm diamond paste with a final step at 0.05 µm colloidal silica. SEM imaging was performed using a Thermo Fisher Apreo FEG SEM. SEM imaging was conducted at 5 kV accelerating voltage and a beam current of 1.6 nA. To study the fine scale substructure, defect content and second phases present in G91 weldments, samples were prepared for scanning transmission electron microscopy (STEM) investigation. To allow for site-specific sample extraction as well as to reduce the magnitude of the ferromagnetic interactions with the beam, samples were prepared using focused ion beam (FIB) milling techniques on a Thermo Fisher Helios 600 30kV dual beam FIB. The ICHAZ was located using a combination of chemical etching and microhardness mapping, as illustrated in Figure 58. Hardness indents were used as fiducial marks to extract specimens and study the local microstructure in various regions of the HAZ. STEM investigations were performed on a Thermo Fisher Tecnai F20 FEG S/TEM operating at 200 kV and spot size 6. Bright field (BF) and annular dark field (ADF) imaging modes were utilized during this study. Tension-tension creep-fatigue testing was conducted on the load-controlled CF testing apparatus. Operation of the CF testing apparatus is described in detail in section d. R-ratios varied from R=0.1-0.12 for the various tests conducted. Figure 59 shows a schematic stress- strain hysteresis loop from a load-controlled creep-fatigue test on G91. The cross weld specimen geometry used for testing is shown previously in Figure 20. In addition, simulated ICHAZ specimens were created using a Gleeble 3500 thermomechanical simulator, using thermal profiles shown in Figure 18. Simulated ICHAZ specimens for CF testing were machined using the sample geometry shown in Figure 20.

43 f.3 Results and Discussion f.3.1 Initial Microstructure A detailed analysis of base metal microstructure is provided in Section e and a brief summary is provided in the following for comparison with weld microstructure . The martensitic transformation and subsequent tempering in CSEFS produces a complex microstructure, which leads to the favorable creep resistance found in CSEFS. Essential to creep strength are the fine subgrains present in CSEFS. Subgrains are created during the martensitic transformation and subsequent tempering, at which point low-angle laths become equiaxed and mobile dislocations rearrange to form low-energy boundaries. Subgrain boundaries migrate and subgrains coarsen during creep and creep-fatigue. Preventing subgrain coarsening is essential to maintaining creep and creep-fatigue strength. [29], [34] Subgrain boundaries are pinned by a fine dispersion of 100-200nm diameter Cr-Mo rich M23C6 precipitates which reside at low-angle boundaries. In addition, fine V-rich carbonitrides exist within the matrix and pin mobile dislocations during elevated temperature deformation. [107] The STEM-DCI DF images in Figure 60 illustrates the microstructure of undeformed FCAW ICHAZ specimens. In Figure 60a the fine substructure is readily apparent, and substructure pinning M23C6 precipitates are observed at low angle boundaries. The fine MX precipitates are observed in Figure 60b, interacting with individual mobile dislocations. f.3.2 Full Weldment Creep-fatigue Testing Creep-fatigue tests were conducted at 650°C with 900s dwell times for CMT and FCAW weldments. For all testing conditions, CMT samples outperformed the FCAW samples. Figure 61 shows the accumulated strain vs. time curves for the FCAW and CMT samples at 100MPa and 85MPa maximum loads and 10 MPa minimum load. For simplicity, only the maximum strain and the time at which the max strain occurred for each cycle is plotted. From this figure it is clear that the CMT samples significantly outperform the FCAW samples. The time to rupture (tr) for the CMT samples is 5x-10x longer than the time to rupture for FCAW. In addition, the minimum creep strain-rate (휀miṅ ) for the CMT samples is an order of magnitude lower at 85 MPa and 100 MPa maximum load test conditions. The CMT welding process results in a reduction in HAZ width compared to the FCAW process. While the peak temperature in the CMT and FCAW welding processes is equivalent, the cooling rate is significantly faster in the CMT welding process, due to the low heat-input from the pulsed arc. Figure 62 shows OM images taken from PWHT and etched CMT and FCAW specimens, illustrating HAZ a width reduction of 40% in the CMT specimens. To explain the increase in CF strength resulting from the CMT welding process it is necessary to examine the fine scale microstructure present in the ICHAZ, where failure occurred for all specimens tested, where fracture occurred for all specimens tested. Examples of failed FCAW and CMT specimens are shown in Figure 63a and Figure 63b, respectively. Both weldments failed via creep void coalescence, as shown in Figure 63c and Figure 63d.

44 Using 5 kV SEM imaging in combination with STEM-DCI, extensive characterization studies were conducted to evaluate changes in precipitate size and distribution, substructure coarsening and defect content in the ICHAZ for the PWHT specimens as well as CF specimens tested at 100 MPa maximum load. Figure 64 shows SEM backscattered electron (BSE) images from the ICHAZ, showing Cr and Mo rich M23C6 precipitates decorating high and low-angle grain boundaries for ICHAZ specimens in the PWHT and post-CF failure conditions. Segmentation of the SEM-BSE images from each weldment made it possible to acquire precipitate size and distribution characteristics. Table 7 shows precipitate characteristics before and after CF deformation.

Table 7: M23C6 characteristics before and after CF deformation for CMT and FCAW weldments.

Nearest Area Equivalent Coarsening Neighbor Sample Fraction Diameter Rate (nm3/s) Distance M C (%) (nm) 23 6 (nm) BM-PWHT 1.96±0.10 102.3±4 N/A 343.6±22.4 FCAW- PWHT 1.86±0.16 105.2±4.4 350.7±35.2 8.96 FCAW- CF 1.74±0.32 145.8±8.0 557.5±60.1 CMT-PWHT 2.1±0.29 111±10 359±25 0.53 CMT-CF 2.19±0.38 126±9 427±58

The equivalent diameter cumulative probability distributions are shown below in Figure 65. M23C6 area fraction for the CMT specimens was found to be 2.1% in the ICHAZ after PWHT while M23C6 area fraction for the FCAW specimen was found to be 1.83% in the ICHAZ after PWHT. In addition, the mean equivalent diameter and average nearest neighbor distance showed a larger increase in the FCAW weldments after CF testing than for the CMT weldments, even though time to rupture (tr) for the CMT weldments was 5x longer than for the FCAW specimens. Figure 65 also shows that the size distribution of M23C6 precipitates in CMT and FCAW ICHAZ after PWHT is almost identical. However, after CF deformation the FCAW weldments show a pronounced shift to larger precipitate sizes in the ICHAZ, indicating that precipitate stability is superior in the CMT weldments. It is important to note that the time to rupture for identical test conditions was 5.6x longer for the CMT specimens, as indicated by the coarsening rates given in Table 7. Dislocation recovery and subgrain coarsening have been cited as mechanisms for monotonic creep deformation in CSEFS. [33], [34], [92] It has also been shown previously that M23C6 precipitates are critical for impeding dislocation recovery and subgrain coarsening. [34] Using zone-axis STEM diffraction contrast imaging it was found that there are significant differences in dislocation density and subgrain characteristics for each weldment before and

45 after CF deformation. STEM-DCI imaging offers many advantages over conventional diffraction- contrast techniques including muted bend contours, which make it possible to observe large fields of view for substructure observations without having to tilt the specimen. In addition,

STEM is advantageous for dislocation density measurements, which require zone-axis imaging to activate multiple diffraction conditions at the same time, allowing for a representative view of total dislocation content. [86] Figure 66 shows STEM-ADF images collected from each weldment after PWHT and after failure under CF loading conditions. It also shows the pronounced increase in subgrain size and precipitate coarsening in the FCAW specimens after deformation. Table 8 shows dislocation density and subgrain mean equivalent diameter measurements determined using the line intercept method outlined in ASTM E112. Dislocation density was measured using zone axis images, using Equation (5) from section e which is described extensively in the included references. [90], [108] Figure 67 shows the zone-axis STEM-BF images used to calculate dislocation density, a significant change in dislocation content between the two welding processes is evident. The CMT specimens have a markedly higher dislocation density before and after CF deformation when compared to the FCAW specimens.

Table 8: Dislocation density and subgrain size for CMT and FCAW weldments after PWHT and CF failure.

Sample Dislocation Density (m-2) Subgrain Size (nm) CMT-PWHT 1.19x1014 ± 1.04 x 1014 581±121 CMT-CF 1.99x1014 ± 1.61x1014 517±76 FCAW-PWHT 5.20 x1013 ± 3.11 x1013 691±76 13 13 FCAW-CF 4.04 x10 ± 3.40x10 1088±480

Dislocation density remains constant in the ICHAZ for the CMT and FCAW specimens before and after CF deformation. However, the dislocation density is greater in the CMT ICHAZ than the FCAW ICHAZ. The finer subgrain size and increased dislocation density in the CMT weldments may contribute to the decreased creep ductility in the CMT specimens. In addition, the increased dislocation density and refined subgrain size in the PWHT CMT specimens as compared to the PWHT FCAW specimens most likely leads to the general increase in HAZ hardness, illustrated in Figure 58. Also, important to note is the subgrain coarsening in the FCAW specimens compared to stable substructure present in the CMT ICHAZ. f.3.3 Simulated ICHAZ Creep-fatigue testing Using the thermal profiles from the ICHAZ during welding, simulated ICHAZ specimens were created using a Gleeble 3500 thermomechanical simulator. These specimens were PWHT

46 at 760°C for 2 hours and tested under identical CF conditions as the full weldments. Prior to CF testing of the Gleeble simulated ICHAZ specimens, STEM specimens were extracted from the simulated ICHAZ specimens to ensure the microstructure was similar to the ICHAZ of the full weldments. Figure 68 shows STEM-ADF images from the Simulated ICHAZ confirming that the microstructure does match the ICHAZ microstructure from the full weldments, which is shown in Figures 66 and 67. Subgrain sizes from the simulated ICHAZ specimens after PWHT were measured and are shown in Table 9 with the subgrain sizes from the full weldment specimens. The similar subgrain size measurements and matching microstructures confirm that the thermal profiles used for the ICHAZ simulation were accurate and that the ICHAZ simulations were successful.

Table 9: Comparison of subgrain sizes between simulated CMT and FCAW ICHAZ specimens and full CMT and FCAW weldments.

Welding Process Sample Subgrain size (nm) Full Weldment 691±76 FCAW Gleeble SIM ICHAZ 857±58 Full Weldment 581±121 CMT Gleeble SIM 540±91 ICHAZ

Figure 69 shows the results from the simulated HAZ CF testing. It shows that the difference in CF performance between the FCAW and CMT simulated ICHAZ samples is negligible at 100 MPa maximum load and the CMT simulated ICHAZ specimen slightly outperforms the FCAW simulated ICHAZ specimen at 85 MPa max load. The tr and 휀푚푖푛̇ for the simulated FCAW ICHAZ specimens are consistent with tr and 휀푚푖푛̇ for the full weldment FCAW specimens. However, the tr and 휀푚푖푛̇ for the CMT specimens is significantly worse than those of the full weldment CMT CF specimens. Thus, there is a discrepancy in CF response between full weldment and Gleeble simulated HAZ specimens, which will be discussed further. The combined results of the Gleeble simulated ICHAZ CF specimens as well as the full weldment CF specimens are shown together in Figure 70. The difference in CF rupture life is negligible between the Gleeble simulated ICHAZ specimens at 100 or 85 MPa. In addition, the simulated ICHAZ rupture lives are comparable to the FCAW weldments at each load, however, the FCAW specimens show slightly lower minimum strain rates. The only specimens that show significantly better CF performance are the CMT full weldment specimens, which exhibit significantly lower minimum strain-rates and an order of magnitude longer rupture lives than the simulated ICHAZ specimens or FCAW full weldments.

47 The low heat-input achieved by using CMT to weld G91 contributes significantly to the CF strength of G91 weldments. These results are consistent with those of Albert et al who showed that CSEFS welded with lower heat-input lead to increased monotonic creep strength. [6] Macroscopically, the decreased heat-input produced a HAZ width that was 40% smaller than the HAZ width for the conventional FCAW weldment. In addition, the decreased heat- input of the CMT welding process results in a cooling rate in the ICHAZ that is 3x faster than the cooling rate of the conventional FCAW welding process. Examination of Figure 64 and Table 7 show that the precipitate characteristics are more favorable for elevated temperature strength in the CMT ICHAZ than in the FCAW ICHAZ. While the equivalent diameter of the M23C6 precipitates is approximately equal for the CMT and FCAW ICHAZ after PWHT, the area fraction is larger and average nearest neighbor distance is smaller for the ICHAZ in the CMT weldment. In addition, the mean equivalent diameter, and average nearest neighbor distance are more stable in the ICHAZ of the CMT weldment, exhibiting less degradation after CF deformation. Previous experiments have shown that increasing precipitate number density and decreasing interparticle spacing, either by alloying or thermomechanical processing can lead to increased monotonic creep strength in CSEFS BM and weldments. [6], [7], [49], [109] The positive effect of M23C6 carbides on elevated temperature strength of CSEFS has been well documented. [7], [29], [32], [33] M23C6 precipitates contribute to creep and creep-fatigue strength by pinning subgrain boundaries (by a Zener pinning mechanism) and preventing subgrain coarsening. It has also been suggested that M23C6 precipitates inhibit mobile dislocation knitting reactions with boundaries. [29], [34] Base metal studies indicate that one recovery mechanism for CSEFS, especially during primary creep is the reduction in dislocation density, some of which must occur by the thermally activated climb and glide of dislocations into subgrain boundaries. [105] By obstructing knitting reactions, the reduction in dislocation density is prevented, further increasing the elevated temperature strength of CSEFS.

The improved M23C6 distribution during CF in the CMT weldments is the direct result of an increase in volume fraction of M23C6 precipitates in the CMT ICHAZ during PWHT. Yu et al have shown previously that an increase in particle density of M23C6 precipitates via thermal processing results in a 5x increase in the creep rupture life of G91 weldments. [7] The increase in creep rupture life of the previously mentioned study was due to a change in pre-weld tempering temperature. The measurements exhibited in Table 7 show that the volume fraction of M23C6 precipitates in the PWHT CMT ICHAZ is increased over the BM and the PWHT FCAW ICHAZ, while the average particle size remains constant, which indicates an increase in precipitate density. f.3.4 Weldment Diffusion Simulations To help explain this increase in volume fraction it is important to make note of the composition of the filler metal of each weldment with comparison to the base metal as well as the width of the HAZ in the FCAW and ICHAZ weldments. The ICHAZ is located on the outer

48 edge of the HAZ, which lies 1.1mm from the fusion boundary (FB) in the CMT weldments and 2.0mm from the FB of the FCAW weldments. Table 4 shows the composition of the solid and flux-cored wires used to weld the CMT and FCAW specimens, respectively. The amount of carbon is higher in the CMT filler wire than in the flux-cored wire or BM. Using DICTRA to model carbon diffusion in the CMT and FCAW weldments it is seen that carbon readily diffuses from the FZ into the HAZ during welding and PWHT. The increase in carbon concentration in the HAZ can lead to more favorable M23C6 size and volume fraction during PWHT and CF, subsequently increasing the CF strength of the CMT ICHAZ. Figure 71 shows results from DICTRA simulations used to model carbon diffusion from the FZ to the ICHAZ of the FCAW and CMT weldments. The model utilizes a simplified specimen geometry and thermal profile. The model is based on a single-phase BCC diffusion couple in which one side starts with the BM composition and one side starts with the filler metal composition. The TCFE9 thermodynamic database and MOBFE4 mobility database were used for the simulation. Since diffusion at the center of the diffusion couple (which is simulating the fusion boundary in this case) is the most important region, a double geometric grid with 60 points was utilized. The width for each diffusion couple was chosen based on the HAZ widths measured using OM, shown in Figure 62. To reduce the simulation complexity, a simplified 5 component Fe-C-Cr-Mo-V composition was utilized, the simplified compositions for the base metal and each filler metal are shown in Table 10. A trapezoidal time integration method was used for the present simulation.

Table 10: Simplified alloy compositions used for DICTRA simulations.

Composition C Cr Mo V Fe (wt%) BM 0.08 8.40 0.92 0.23 Bal FCAW WM 0.096 8.73 1.03 0.19 Bal CMT WM 0.106 8.92 0.99 0.22 Bal

It is important to note that these models do not account for the effect of dilution during welding or reheating caused by multipass welding. Dilution should be greater in the higher heat-input FCAW process, resulting in even less driving force for carbon diffusion after welding in FCAW than CMT. The resulting diffusion couple is run through the same thermal profiles used to simulate the ICHAZ specimens, which were measured during welding of the CMT specimens, and extracted from literature for the FCAW specimens. After simulating the welding profiles, the diffusion couple is also run through PWHT (2 hours at 760C) and thermal aging at the CF test temperature of 650C. Figure 71 shows that the amount of carbon is greater in the CMT ICHAZ than in the FCAW ICHAZ after PWHT. Increased carbon content accounts for the reduced increase in mean

49 equivalent diameter and interparticle spacing in the CMT specimens by allowing for the nucleation of new fine M23C6 precipitates while the existing precipitates coarsen during PWHT and aging. The refined distribution of M23C6 precipitates inhibits subgrain coarsening and knitting reactions, leading to an increase in elevated temperature strength. It is important to note the similarity of the rupture life and minimum strain rate of the ICHAZ simulated CF tests and the FCAW specimens, shown in Figure 70. The ICHAZ simulated specimens do not experience any carbon diffusion and are made up entirely of BM composition material subjected to the thermal history of the ICHAZ. Without the increase in carbon to nucleate new M23C6 precipitates the simulated weldments do not exhibit increased resistance to CF deformation. There is no significant difference in CF performance of the simulated CMT weldments or simulated FCAW weldments at 100 or 85 MPa maximum dwell loads. This shows that the decreased heat-input (and resulting decrease in HAZ width) as well as the increased carbon content in the weld metal are both necessary to enhance the CF performance of the CMT weldments.

Combining the M23C6 precipitation characteristics from PWHT and post-CF deformation in Table 7 with the cumulative probability curves for M23C6 diameter in Figure 65 shows that during PWHT and CF in the CMT specimens there is increased nucleation of M23C6 precipitates and coarsening of existing precipitates while in the FCAW ICHAZ there is less M23C6 nucleation during PWHT and little to no nucleation during CF deformation. This observation agrees with the increased carbon content indicated by the DICTRA simulations. While the FCAW full weldment DICTRA simulations indicate that there is some carbon diffusion during PWHT, the experimental results indicate otherwise. The fact that the volume fractions do not increase during CF and also that the rupture times and minimum strain rates are very similar for the full weldments and simulated ICHAZ specimens (where there is no gradient in carbon content) indicates that there is minimal net carbon transport occurring in the FCAW specimens. It is important to note that the FCAW process involves considerable dilution during welding compared to lower heat-input processes. Increased dilution would move excess carbon from the FZ to HAZ during welding and decrease the solute gradient and driving force for carbon diffusion during PWHT and CF deformation. Also, the FCAW wire is flux-cored, as compared to the solid filler wire used in the CMT welding process. It is possible that some of the carbon in the flux is lost during welding and does not make it into the FZ during the welding process. Conversely, the CMT filler wire is solid and is more likely to result in matching filler wire and FZ compositions. f.4 Local Strain Accumulation in HAZ Since cross-weld specimens contain a microstructure gradient and thus property gradient along the tensile direction, the local deformation (as opposed to the total deformation over the gauge section) is essential to study the CF damage evolution. A high-temperature DIC technique was developed to observe the local strain distribution in-situ during CF testing. The

50 max load used during the high-temperature DIC experiment was 150 MPa. Due to equipment limitations, the total testing duration was limited to 10 hours. Testing at the lower max. loads (e.g., 100 and 85 MPa) did not result in fracture of cross-weld specimens. High-temperature DIC results for a FCAW cross-weld specimen are shown in Figure 72, where the boundaries of weld metal (WM), heat-affected zone (HAZ), and base metal (BM) are marked. As shown in this figure, strain rapidly accumulated in the HAZ even after the first two cycles. The strain continued to accumulate in the HAZ, and after the last cycle of the test (cycle #29), the highest strain was found in the HAZ adjacent to the BM. This HAZ region is expected to be ICHAZ/FGHAZ. It is noted that the test was interrupted after cycle #29 due to equipment limitation, and the cross-weld specimen was not fractured. A comparison of the local strain in HAZ as a function of time for FCAW and CMT cross- weld specimens is shown in Figure 73. Due to the extremely large volume of data generated by the DIC camera, two cycles were measured and analyzed for FCAW specimen; four cycles for CMT specimen. The actually measured data are plotted as solid lines whereas the data between the recording cycles are extrapolated as dashed lines. Similar to that shown in Figure 61, the CMT specimen fractured at a much longer time (7.5 hours) than the FCAW specimen (2 hours). The steady-state strain rate is also much lower for CMT than FCAW. The local HAZ strain at fracture is about 50% for CMT. For FCAW, it is noted that the final fracture strain was not determined as the DIC recording was not done for the cycle during which the specimen failed. To better quantify the strain accumulation and failure location, the strain field measured at the cycle closest to the final failure for FCAW vs. CMT was compared in Figure 74. This strain field was superimposed on a “un-deformed” reference image prior to CF testing. As shown in this figure, the location with highest local strain and failure measured on the un-deformed image was at a distance of 1.95 mm and 1.04 mm away from the fusion boundary for the FCAW and CMT specimens, respectively. These distances are fairly consistent with the HAZ widths shown in Figure 62. Combining with the post-test optical microcopy, the failure was determined to occur in the ICHAZ/FGHAZ for both CMT and FCAW cross-weld specimens. f.5 Summary and Conclusions for G91 Weld Studies The current study uses a load-controlled creep-fatigue testing apparatus to study the effect of low heat-input welding on CF properties of CSEFS. Using conventional FCAW and a non-conventional CMT low heat-input process it was found that CMT weldments out performed FCAW under CF conditions. The following other conclusions were made:  Cold metal transfer welding creates weldments with low heat-inputs compared to other arc welding processes. Welds from this study were produced using a heat-input of 0.67 kJ/mm, which is comparable to EBW or LW heat inputs documented in other studies. [6]  CMT specimens exhibit an order of magnitude lower minimum strain rate than FCAW

51 specimens under load-controlled CF conditions. The time-to-rupture is at least 3.75x longer for the CMT specimens in all test conditions examined. However, failure occurs in the ICHAZ for all specimens, regardless of welding process.  The CMT low heat-input welding process produces a HAZ which is 40% narrower than the FCAW HAZ. The cooling rate in the ICHAZ for the CMT specimens is 3x faster than the FCAW process.

 Previous studies suggest that M23C6 precipitates are critical for retaining elevated temperature strength in CSEFS. The particle size is almost identical for CMT and FCAW after PWHT, however, the area fraction of M23C6 is higher in the CMT specimens. After CF failure, the M23C6 particles in the ICHAZ of the FCAW specimens show more coarsening, and larger interparticle spacing when compared to the CMT specimens. Correspondingly, CMT weldments also exhibit greater substructure stability than FCAW weldments.  Simulated CMT and FCAW ICHAZ bulk specimens are also tested. Results from the simulated specimens were comparable to the FCAW full weldments, with only the CMT full weldments exhibiting enhanced CF strength.  DICTRA simulations show that carbon diffusion from the FZ to the ICHAZ contributes to the increase in area fraction of M23C6 and refined particle distribution in CMT after PWHT and CF deformation. The CMT weldment’s carbon enriched filler metal and narrow HAZ are both necessary to produce the favorable M23C6 distribution.

g. Overall Summary, Conclusions and Future Work g.1 G91 Base Metal Creep-fatigue Studies Section e focuses on the base metal CF behavior of the ferritic-martensitic alloy G91. Initial microstructural observations found that after normalizing and tempering, G91 exhibits a tempered martensitic microstructure with a refined subgrain network stabilized by 100nm- 200nm carbides and a high dislocation density. A systematic CF study was performed on tempered G91 base metal, examining microstructural evolution with creep-fatigue damage as well as the effect of loading parameters on load-controlled creep-fatigue deformation. This study is unique in that a load-controlled CF testing apparatus was designed and built to conduct load-controlled, creep-dominated CF testing, which is an uncommon testing method for these materials but allows for creep-dominated CF testing and the observation of anelastic backflow at minimum load. Load-controlled creep-fatigue strain-time curves strongly resemble monotonic creep strain-time curves for G91 specimens, with significant primary and tertiary transients. Microstructural observations on crept specimens reveal that the large primary transient is due to a significant decrease in dislocation density at the onset of CF testing while the tertiary transient is related to subgrain coarsening and creep-void formation and coalescence.

52 Interrupted CF experiments allow for the study of microstructural evolution as a function of CF deformation. It was found that subgrains coarsen continuously as a function of CF strain, however, this coarsening is counteracted by the significant decrease in mobile dislocation density that occurs during CF testing, some of which must be due to the rearrangement of mobile dislocations into low-energy subgrain boundaries. Eventually, subgrain coarsening rates exceed the rate of dislocation wall formation and the subgrains coarsen rapidly until failure.

The main obstacle to subgrain coarsening is the pinning of LAGBs by M23C6 precipitates, which is also the main strengthening mechanism in monotonic creep deformation of G91. [29] Comparing CF stress exponents determined in the present study to monotonic creep exponents from literature support the idea that the same mechanisms govern CF and monotonic creep deformation. However, strain rates are significantly higher and rupture lives are significantly shorter for CF compared monotonic creep. Correspondingly, subgrain coarsening rates are higher for CF, indicating that the barriers to subgrain migration in G91 are less effective under CF conditions. The accelerated microstructural degradation under CF conditions is closely related to the anelastic backflow that occurs under CF loading conditions, which is due to the forward and backward bowing of subgrain boundaries during CF deformation. Studies concerning anelastic backflow show that anelasticity increases with minimum load and maximum load dwell times, as well as increasing maximum loads. It was determined that to rationalize cycles to failure, forward creep strain and anelastic backflow must be considered to calculate a net forward flow of strain. Anelastic backflow may cause the accelerated subgrain coarsening and reduced rupture lives exhibited during CF deformation.

Future work for G91 base metal CF should include a systematic study of M23C6 coarsening rates for CF of G91 specimens to compare the effect CF loading and anelastic backflow on particle coarsening compared to monotonic creep particle coarsening rates, since the M23C6 particles are the primary obstacles to subgrain migration. In addition, modeling efforts are necessary to quantify the forward bowing and relaxation of subgrain boundaries during cyclic loading. g.2 G91 Low heat-input Welding and Creep-fatigue Studies Creep-fatigue deformation is especially detrimental to G91 weldments, with failure typically occurring in the outer region of the HAZ, specifically in the FGHAZ or ICHAZ. [18] T his type of failure is common for CSEFS welded components subjected to elevated temperature deformation. In the current study, CF performance of conventional FCAW weldments is compared to non-conventional low heat-input CMT weldments. Under multiple loading conditions CMT specimens significantly outperform the FCAW specimens, with rupture times at least 3.75x longer than the FCAW specimens and an order of magnitude lower 휀푚푖푛̇ than FCAW full weldment specimens. The increase in CF strength in the CMT weldments is the result of a refined distribution of M23C6 precipitates, which are key to creep-fatigue strength in G91. The refined M23C6

53 distribution more effectively pins subgrain boundaries, which was reflected in the refined subgrain size and increased dislocation density present in the CMT specimens after PWHT and CF deformation compared to the FCAW specimens.

In addition, the area fraction of M23C6 is 0.2% higher in the CMT weldments after PWHT. The increase in precipitate volume fraction and refinement in precipitate distribution is due to a combination of low heat-input welding, which produces a 40% narrower HAZ in the CMT specimens as well as the carbon enriched filler metal utilized in the CMT welding process. After CF rupture, the CMT specimens show a more refined precipitate distribution compared to FCAW, with smaller mean equivalent diameters and reduced nearest neighbor distances, despite being at elevated temperatures under CF deformation conditions for significantly longer periods of time. Carbon diffusion from the FZ to the ICHAZ was confirmed with DICTRA simulations. Simulated ICHAZ specimens which were subjected to ICHAZ thermal cycles but contained G91 BM composition exhibit CF performance similar to FCAW full weldments, indicating that carbon enrichment from the FZ is necessary to increase CF rupture strength. The carbon enrichment from the FZ to the ICHAZ results in a refined distribution of M23C6 precipitates and a greater area fraction which increases substructure hardening and mitigates the type IV cracking phenomenon. g.3 Overall Conclusions The principal goal of the analysis was to provide mechanistic details that correlate processing and CF loading parameters to microstructure and elevated temperature deformation in high-alloy G91 base metal and welded CSEFS. These alloys are used extensively in the power generation industry and increased understanding of deformation in these alloys is critical to increased energy production efficiency. G91 creep-fatigue research conducted on base metal plate provides a more accurate description of deformation in CSEFS under cyclic loading conditions. The studies regarding microstructural evolution and the effect of loading parameters on CF life provide new insight into load-controlled CF deformation of CSEFS. In addition, the research concerning low-heat input arc welding of G91 components provides evidence of type IV cracking mitigation in CSEFS, which could improve component lifetime for G91 components, which are usually limited by weldment failure. CSEFS are a workhorse material for the power generation industry, the present body of work provides new insights into how these materials respond to processing conditions and elevated temperature deformation, which can be used to produce improved creep-strength enhanced ferritic- martensitic steel weldments or more accurate deformation models for CSEFS components.

h. List of publications Theses:  T. K. Payton, “On the Improvement of Creep-fatigue Behavior of Grade 91 Weldments,” The Ohio State University, M.S. Thesis in Welding Engineering, July 2017. http://rave.ohiolink.edu/etdc/view?acc_num=osu150185172211282

54  H.C. Whitt, “Creep and Creep-fatigue Deformation Studies in 22V and G91 Creep- strength Enhanced Ferritic Steels” The Ohio State University, Ph.D. Dissertation in Materials Science and Engineering. May 2019.

Journal Papers:  H. C. Whitt, T. K. Payton, W. Zhang, and M. J. Mills, “Analysis of Creep-Fatigue Deformation and Microstructural Evolution in 9Cr-1MoV Steel,” manuscript in preparation, May 2019.  H. C. Whitt, T. K. Payton, W. Zhang, and M. J. Mills, “Creep-Fatigue Properties of 9Cr- 1MoV Low Heat Input Weldments,” manuscript in preparation, May. 2019.  K. Zhang, H. C. Whitt, W. Zhang, and M. J. Mills, “Strain Localization in Heat-Affected Zone of Grade 91 Weld during Creep-Fatigue Cycling,” manuscript in preparation, May. 2019.  T. Mukherjee, J. S. Zuback, W. Zhang and T. DebRoy, “Residual stresses and distortion in additively manufactured compositionally graded transition joints,” Computational Materials Science, Vol. 143, pp. 325-337 (2018). https://doi.org/10.1016/j.commatsci.2017.11.026

Conference presentations:  H. C. Whitt, T. K. Payton, M. J. Mills, W. Zhang, Y. Wang, Z. Feng, “Creep-Fatigue Interactions in 9Cr-1MoV Steel and Weldments,” DOE Advanced Reactor Technologies: Advanced Materials Program Review Meeting, Idaho Falls, Idaho, 14-Jul-15.  H. C. Whitt, T. Payton, Y. Wang, W. Zhang, M.J. Mills, “"Creep-Fatigue Interactions in 9Cr-1MoV Steel and Welds," 2015 Materials Science and Technology (MS&T), Columbus, Ohio, October 5th, 2015.  T. K. Payton, H. C. Whitt, W. Zhang, M. J. Mills, Y. Wang, “Creep-Fatigue Performance of 9Cr-1MoV Weldments,” 2015 Materials Science and Technology (MS&T), Columbus, Ohio, October 5th, 2015.  H. C. Whitt, T. K. Payton, Y. Wang, W. Zhang, M. J. Mills, “Creep-Fatigue Interactions in 9Cr-1MoV Steel and Welds,” FabTech / AWS Annual Convention 2015, Chicago, IL, 10- Nov-15.  H. C. Whitt, T. K. Payton, Y. Wang, W. Zhang, M. J. Mills, “Creep-Fatigue Interactions in 9Cr-1MoV Steel and Welds,” TMS 2016, Nashville, TN, 18-Feb-16.  H. C. Whitt, T. K. Payton, W. Zhang, M. J. Mills, “Creep-Fatigue Interactions in 9Cr-1MoV Steel and Weldments,” GE Global Research Student Research Summit, Niskayuna, NY, 8/1/2016.

55  T. K. Payton, H. C. Whitt, W. Zhang, M. J. Mills, Y. Wang, “Effect of Low Heat Input on Creep Strength of 9Cr-1MoV Weldments,” FabTech/AWS Annual Convention 2016, Las Vegas, NV, 11/4/2016.  H. C. Whitt, T. K. Payton, W. Zhang, M. J. Mills, “Creep-Fatigue Deformation of 9Cr- 1MoV Steel and Weldments, “ TMS 2017, Materials and Fuels for the Current and Advanced Nuclear Reactors VI, San Diego, CA, 2-Mar-17.  H. C. Whitt, T. K. Payton, W. Zhang, M. J. Mills, “Creep-Fatigue Interactions in 9Cr-1MoV Steel and Weldments,” Gordon Research Conference - Physical , Biddeford, ME, Jul-17.  T. K. Payton, H. C. Whitt, M. J. Mills, W. Zhang, Y. Wang, “Effect of Low Heat Input on Creep Strength of 9Cr-1MoV Weldments,” FabTech/AWS Annual Convention 2017, Chicago, IL, Nov-17.  H. C. Whitt, T. K. Payton, W. Zhang, M. J. Mills, “Characterization of Creep-Fatigue Deformation in 9Cr-1MoV Steel and weldments,” TMS, Phoenix AZ. Mar-2018.  H. C. Whitt, T. K. Payton, W. Zhang, M. J. Mills, “Characterization of Creep-Fatigue Deformation in 9Cr-1MoV Steel and weldments,” MS&T, Columbus, OH. Oct-2018.  H. C. Whitt, T. K. Payton, W. Zhang, M. J. Mills, “Creep-fatigue deformation in 9Cr-1MoV Steel,” TMS, San Antonio, TX. Mar-2019.

i. Acknowledgements This project is supported by U.S. Department of Energy, Nuclear Energy University Program (NEUP). The authors are grateful to Drs. Sam Sham and Meimei Li of Argonne National Laboratory for serving as the technical program managers and for providing valuable inputs to the research. In addition, Mr. Joe Bundy, Hobart Brothers Company is acknowledged for supporting welding of Grade 91 plates including supplying base metal, filler wire and baseline welding parameters.

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62

Figure Captions

Figure 1: Evolution of CSEFS over the last 8 decades, with corresponding 105 hr rupture stress.[5] Figure 2: Microstructure representative of 9-12% chromium CSEFS. (a) Optical micrograph showing PAGBs. (b) High-angle annular dark-field (HAADF) scanning transmission electron micrograph (STEM) showing significant dislocation content inside subgrains and precipitates at subgrain boundaries.[54],[55] Figure 3: Peak stress vs. cycle number illustrating cyclic softening in G91. (a) CF experiments performed on G91 at ANL (hold time 60s in tension and compression). (b) Examples of cyclic softening exhibited by P92 specimens with various dwell times in tension and compression. [16], [38] Figure 4: Plots illustrating the effect of hold time on cycles to failure for (a) hold time under a tensile stress and (b) hold time under a compressive stress. [15]

Figure 5: STEM images illustrating: (a) M23C6 precipitates at LAGB and PAGB. (b) Substructure present after tempering and CF deformation. Interactions between MX carbonitrides and mobile dislocations (c) before and (d) after CF deformation. [38] Figure 6: Interaction diagrams for G91 experiments using the (a) TFA and (b) DEM. (a) shows the failure envelope determined by ASME. [15] Figure 7: Comparison of CF life prediction to experimental lives for (a) the TFA, (b) the DEM and (c) the modified DEM. [15] Figure 8: (a) Schematic example of a total load-controlled CF hystersis, exhibiting anelastic backflow of strain at minimum load. (b) Example of a partially load controlled CF test, which is strain-controlled during loading, not allowing for the accumulation of anelastic strain. (c) Schematic example of a strain-controlled CF hystreresis, not allowing anleastic backflow during minimum load. [39] Figure 9: Displacement vs. time curves from Sawada et al, showing the effect of microstructure on anelastic backflow upon abrupt stress drops. (a) CSEFS with substructure and M23C6 substructure stabilizing precipitates. (b) CSEFS with substructure but no substructure stabilizing precipitates. (c) CSEFS with no substructure. (d) CSEFS with fine distribution of MX carbonitrides stabilizing lath boundaries. [48] Figure 10: Schematic representation of HAZ in G91, correlating HAZ regions with peak temperatures during welding. [18], [53] Figure 11: EBSD IPF map from standard pre-weld heat treatment and (b) the corresponding IQ map showing ferrite formation. (c) and (d) IPF and IQ map from low-temperature heat

63 treatment showing no ferrite formation. [12] Figure 12: a) Comparison of creep curves from experimental and simulated creep tests on P122 weldments. (b) Experimental results from tensile tests on simulated HAZ specimens from P122 weldment. [13] Figure 13: a) Creep rupture plot illustrating rupture life of simulated FGHAZ specimens compared to welds and base metal. (b) Relationship between heat input (noting different welding processes), HAZ width and weld geometry on rupture life. [6] Figure 14: Creep rupture plots illustrating creep strength of G91 compared to P92 and MARBN steels. (b) Dislocation density and lath width as a function of time under creep in G91 and W containing steel. [24],[63] Figure 15: (a) Schematic showing operation of FCAW welding process. (b) High speed camera capturing short circuit arc droplet during CMT process. [74]–[76] Figure 16: Weldment geometry used for welding G91 CMT and FCAW weldments. Figure 17: Weld Sequence for FCAW and CMT weldments. [77] Figure 18: Welding thermal profiles measured in the ICHAZ for a single pass. Figure 19: Experimental setup inside Gleeble for ICHAZ simulations. Figure 20: G91 base metal and weldment specimen geometry (dimensions in inches). Figure 21: Schematic representation of the creep-fatigue testing apparatus. Figure 22: CF stress-strain hysteresis (left) and loading waveform (right) characteristic of the tests performed in this study. Figure 23: (A) Schematic setup for SE detector. (B) Schematic setup for BSE detector in an SEM. [82] Figure 24: (a) Schematic setup for (a) EBSD and (b) TKD. [83] Figure 25: Schematic detector setup for STEM imaging, showing different angular ranges used for different imaging modes. [87]

Figure 26: (a) Schematic sample with drilled hole at =0 (b) Schematic sample with drilled hole at =25. (c) STEM-HAADF image at =0. (d) STEM-HAADF image at =25.

Figure 27: (Top) BSE image and (Bottom) corresponding SE image, showing M23C6 precipitates at high-angle and low-angle grain boundaries.

64 Figure 28: (Left) STEM-ADF and (right) STEM-BF images showing subgrains, with M23C6 particles decorating subgrain boundaries and high dislocation density. Figure 29: STEM-BF image, showing dislocations bowing around MX vanadium-rich carbonitrides. Figure 30: Time to rupture for monotonic creep (data adapted from Cerri et al and Tabuchi et al) compared to time to rupture for load-controlled CF data collected in the present study. [32], [51] Figure 31: [111] Zone-axis image from a ruptured G91 CF specimen. An overlaid grid with dislocation intersections illustrates the method used to calculate dislocation density. Figure 32: Cumulative strain-per-cycle vs cycle number for a creep-fatigue test. Interrupted tests used to evaluate microstructural evolution are designated by marker circles overlaid on the CF curve. Figure 33: Measured subgrain size as a function of creep-fatigue strain accumulation. For reference, strain rate vs cycle number curve is also presented. Figure 34: (A) STEM-DCI illustrating subgrain coarsening as a result of load-controlled CF deformation. (B) Low-magnification STEM-BF image, emphasizing subgrain coarsening after rupture. Figure 35: (A) Dislocation density as a function of CF cycle number, shown on a linear plot emphasizing significant decrease at the onset of CF deformation. (B) Dislocation density as a function of CF cycle number, shown on a logarithmic plot, emphasizing logarithmic relationship between dislocation density and cycle number. Figure 36: Examples of zone-axis STEM images used to calculate dislocation density for microstructural evolution. Figure 37: EBSD scans taken from interrupted specimens, showing a reduction in low-angle boundaries as a function of CF strain.

Figure 38: Low angle boundary content (less than 15 misorientation) as a function of CF cycle number. Figure 39: Comparison of subgrain coarsening during CF deformation in the present study to subgrain coarsening during monotonic creep deformation studies from existing literature. [37],[113]

Figure 40: Examples of low-angle boundaries being pinned by M23C6 precipitates. Figure 41: Cumulative probability for precipitate size in G91 specimens at different stages of CF strain accumulation.

Figure 42: Strain-rate vs maximum applied stress for creep-fatigue data collected at 600C and 650C with a 900s dwell time, creep data from literature is shown for comparison. [32], [100],

65 [101] Figure 43: Stress-strain hysteresis loops from the primary, secondary and tertiary regimes of a creep-fatigue test with 85 MPa max load and 900s dwell time at 650C. Figure 44: Creep strain and anelastic strain per cycle as a function of cycle number for the CF curve shown in Figure 32. Figure 45: Time at maximum load and time at minimum load vs. anelastic strain for specimens tested with the same maximum and minimum dwell loads. Figure 46: Effect of creep load on anelastic strain and rupture life for CF specimens with matching maximum and minimum load dwell time. Figure 47: Heterogeneous distribution of dislocation content which is characteristic of CSEFS under viscoplastic deformation conditions. (A) BF STEM-DCI showing heterogeneous dislocation content in adjacent subgrains. (B) TKD IPF Map showing low misorientation between subgrains with dissimilar dislocation densities. (C) GROD map illustrating difference in dislocation density (in good agreement with A). Figure 48: (A) STEM-ADF image showing an edge-on subgrain boundary and holes drilled for EDS drift-correction. (B) EDS map showing Fe deficiency, (C) EDS map showing Cr enrichment, (D) EDS map showing Mo enrichment. Figure 49: (A) Schematic stress-strain hysteresis showing the subgrain boundary-solute interaction state at different points during the loading cycle. (B) Solute drag occurring during monotonic creep and the later stages of the dwell period during CF deformation. C) Anelastic backflow allowing boundary dislocations to break free from solute atmosphere. (D) Subgrain migrates without solute atmosphere until boundary dislocations recollect solute. Figure 50: Strain vs. cycle number for CF tests with different maximum load dwell times, at 650C and 85MPa maximum load. Figure 51: Anelastic backflow per cycle for specimens with varied maximum load dwell time. Figure 52: Creep strain per cycle for CF specimens with varied maximum load dwell time. Figure 53: Average anelastic backflow, creep strain, and net forward strain per cycle as a function of dwell time. Figure 54: Net strain per cycle vs. cycles to failure for CF specimens with varied maximum load dwell time. Figure 55: (A) Fracture surface of a load-controlled CF specimen, showing evidence of ductile fracture and microvoids. (B) Polished fracture surface, showing coalescence of creep voids. (C) Wide field of view (FOV) cross-section of a G91 CF specimen after rupture, showing creep voids and coalescence at the fracture surface. (C) Narrow FOV cross-section, showing fine creep voids at GBs and precipitates.

66 Figure 56: G91 normalized and tempered microstructure, left: optical micrograph showing PAGB and martensite. Right: STEM-ADF diffraction contrast image showing high dislocation density, fine subgrains and second phase particles at subgrain boundaries. Figure 57: Chemically etched HAZ from a G91 weldment (CMT-as welded). Figure 58: (a) Chemically etched and hardness indented CMT specimen illustrating method for finding various regions in the HAZ. (b) Microhardness measurements, showing HAZ traverses from the FCAW and CMT weldments. Figure 59: Schematic stress-strain hysteresis loop from a load-controlled CF test on G91. Figure 60: STEM-DCI images from the FCAW ICHAZ: (a) ADF image illustrating fine subgrains and M23C6 precipitates at subgrain boundaries. (b) BF image illustrating significant dislocation content within subgrains, M23C6 at subgrain boundaries and MX carbonitrides interacting with mobile dislocations in subgrain boundaries. Figure 61: Creep-fatigue performance comparison for CMT and FCAW weldments. Figure 62: Optical micrographs showing HAZ width reduction in the CMT specimens when compared to the FCAW specimens. Figure 63: Cross-section showing ICHAZ failure in (a) FCAW and (b) CMT weldments. Creep-void coalescence at the fracture surface of (c) CMT and (d) FCAW weldments. Figure 64: SEM images illustrating precipitate distributions for G91 weldments before and after CF deformation.

Figure 65: Cumulative probability for M23C6 equivalent diameter for weldments before and after CF deformation. Figure 66: STEM-ADF images showing low-angle boundary (1º-5º misorientation) subgrain structure in CMT and FCAW weldments

67 Figure 67: Dislocation density after PWHT (Left) and CF-rupture (Right) for CMT weldments (Top) and FCAW weldments (Bottom). Figure 68: Representative microstructures from simulated ICHAZ specimens, confirming successful ICHAZ microstructure simulations. Figure 69: CF results from simulated HAZ CF tests performed at 100 MPa and 85 MPa max dwell loads. Figure 70: Combined full weldment and simulated ICHAZ CF data. Figure 71: DICTRA simulations showing carbon diffusion from the FZ to the ICHAZ during welding and PWHT of FCAW (Top) and CMT (bottom) weldments. Figure 72: Distribution of longitudinal strain in a cross-weld specimen of FCAW, which shows a significant strain accumulation in the ICHAZ. Strain map was measured using digital image correlation. Testing temperature = 650 C, peak stress = 100 MPa and dwell time = 900 s. Figure 73: Comparison of local strain in HAZ for FCAW versus CMT cross-weld specimens tested at 150 MPa. Local strain measured by high-temperature DIC technique. Figure 74: Distribution of longitudinal strain prior to failure in cross-weld specimens of (a) FCAW and (b) CMT measured by DIC. Testing temperature = 650 °C, peak stress = 150 MPa and dwell time = 900 s.

68

Figure 1: Evolution of CSEFS over the last 8 decades, with corresponding 105 hr rupture stress. [5]

(a) (b)

Figure 2: Microstructure representative of 9-12% chromium CSEFS. (a) Optical micrograph showing PAGBs. (b) High-angle annular dark-field (HAADF) scanning transmission electron micrograph (STEM) showing significant dislocation content inside subgrains and precipitates at subgrain boundaries. [54],[55]

(a) (b)

Figure 3: Peak stress vs. cycle number illustrating cyclic softening in G91. (a) CF experiments performed on G91 at ANL (hold time 60s in tension and compression). (b) Examples of cyclic softening exhibited by P92 specimens with various dwell times in tension and compression. [16], [38]

Figure 4: Plots illustrating the effect of hold time on cycles to failure for (a) hold time under a tensile stress, and (b) hold time under a compressive stress. [15]

Figure 5: STEM images illustrating: (a) M23C6 precipitates at LAGB and PAGB. (b) Substructure present after tempering and CF deformation. Interactions between MX carbonitrides and mobile dislocations (c) before and (d) after CF deformation. [38]

(a) (b)

Figure 6: Interaction diagrams for G91 experiments using the (a) TFA and (b) DEM. (a) shows the failure envelope determined by ASME. [15]

(a) (b)

(c)

Figure 7: Comparison of CF life prediction to experimental lives for (a) the TFA, (b) the DEM, and (c) the modified DEM. [15]

a b

c

Figure 8: (a) Schematic example of a total load-controlled CF hystersis, exhibiting anelastic backflow of strain at minimum load. (b) Example of a partially load controlled CF test, which is strain-controlled during loading, not allowing for the accumulation of anelastic strain. (c) Schematic example of a strain-controlled CF hystreresis, not allowing anleastic backflow during minimum load. [39]

a b

c d

Figure 9: Displacement vs. time curves from Sawada et al, showing the effect of microstructure on anelastic backflow upon abrupt stress drops. (a) CSEFS with substructure and M23C6 substructure stabilizing precipitates. (b) CSEFS with substructure but no substructure stabilizing precipitates. (c) CSEFS with no substructure. (d) CSEFS with fine distribution of MX carbonitrides stabilizing lath boundaries. [48]

Figure 10: Schematic representation of HAZ in G91, correlating HAZ regions with peak temperatures during welding. [18], [53]

Figure 11: EBSD IPF map from standard pre-weld heat treatment and (b) the corresponding IQ map showing ferrite formation. (c) and (d) IPF and IQ map from low-temperature heat treatment showing no ferrite formation. [12]

(a) (b)

Figure 12: (a) Comparison of creep curves from experimental and simulated creep tests on P122 weldments. (b) Experimental results from tensile tests on simulated HAZ specimens from P122 weldment. [13]

(a)

(b)

Figure 13: (a) Creep rupture plot illustrating rupture life of simulated FGHAZ specimens compared to welds and base metal. (b) Relationship between heat input (noting the different welding processes), HAZ width and weld geometry on rupture life. [6]

(a)

(b)

Figure 14: Creep rupture plots illustrating creep strength of G91 compared to P92 and MARBN steels. (b) Dislocation density and lath width as a function of time under creep in G91 and W containing steel. [24],[63]

a

b

Figure 15: (a) Schematic showing operation of FCAW welding process. (b) High speed camera capturing short circuit arc droplet during CMT process. [74]–[76]

20º

FZ BM BM HAZ 19.05 mm

25.4 mm

Figure 16: Weldment geometry used for welding G91 CMT and FCAW weldments.

Figure 17: Weld Sequence for FCAW and CMT weldments. [77]

Figure 18: Welding thermal profiles measured in the ICHAZ for a single pass.

Figure 19: Experimental setup inside Gleeble for ICHAZ simulations.

Figure 20: G91 base metal and weldment specimen geometry (dimensions in inches).

Figure 21: Schematic representation of the creep-fatigue testing apparatus.

Figure 22: CF stress-strain hysteresis (left) and loading waveform (right) characteristic of the tests performed in this study.

Figure 23: (A) Schematic setup for SE detector. (B) Schematic setup for BSE detector in an SEM. [82]

Figure 24: Schematic setup for (a) EBSD and (b) TKD. [83]

Figure 25: Schematic detector setup for STEM imaging, showing different angular ranges used for different imaging modes. [87]

Figure 26: (A) Schematic sample with drilled hole at =0. (B) Schematic sample with drilled hole at =25. (C) STEM-HAADF image at =0. (D) STEM-HAADF image at =25.

Figure 27: (Top) BSE image and (Bottom) corresponding SE image, showing M23C6 precipitates at high-angle and low-angle grain boundaries.

Figure 28: (Left) STEM-ADF and (right) STEM-BF images showing subgrains, with M23C6 particles decorating subgrain boundaries and high dislocation density.

100 nm

Figure 29: STEM-BF image, showing dislocations bowing around MX vanadium-rich carbonitrides.

Figure 30: Time to rupture for monotonic creep (data adapted from Cerri et al and Tabuchi et al) compared to time to rupture for load-controlled CF data collected in the present study. [32], [51]

Figure 31: [111] Zone-axis image from a ruptured G91 CF specimen. An overlaid grid with dislocation intersections illustrates the method used to calculate dislocation density.

N = 1654

N =451 N =26 N = 1000 N =0

Figure 32: Cumulative strain-per-cycle vs cycle number for a creep-fatigue test. Interrupted tests used to evaluate microstructural evolution are designated by marker circles overlaid on the CF curve.

Figure 33: Measured subgrain size as a function of creep-fatigue strain accumulation. For reference, strain rate vs cycle number curve is also presented.

(a) BM N=26

N-451 Post-Failure

1 µm

(b) Post-failure

2 µm

Figure 34: (A) STEM-DCI illustrating subgrain coarsening as a result of load-controlled CF deformation. (B) Low-magnification STEM-BF image, emphasizing subgrain coarsening after rupture.

Figure 35: (A) Dislocation density as a function of CF cycle number, shown on a linear plot emphasizing significant decrease at the onset of CF deformation. (B) Dislocation density as a function of CF cycle number, shown on a logarithmic plot, emphasizing logarithmic relationship between dislocation density and cycle number.

BM N =26

200 nm 200 nm N =451 Rupture

500 nm 500 nm

Figure 36: Examples of zone-axis STEM images used to calculate dislocation density for microstructural evolution.

N=0 N=26

15 m 15 m

N=451 N=1654 - rupture

Misorientation 1º-5º 5º-15º 15º-65 15 m

Figure 37: EBSD scans taken from interrupted specimens, showing a reduction in low-angle boundaries as a function of CF strain

Figure 38: Low angle boundary content (less than 15 misorientation) as a function of CF cycle number.

Figure 39: Comparison of subgrain coarsening during CF deformation in the present study to subgrain coarsening during monotonic creep deformation studies from existing literature. [37],[113]

Figure 40: Examples of low-angle boundaries being pinned by M23C6 precipitates.

Figure 41: Cumulative probability for precipitate size in G91 specimens at different stages of CF strain accumulation.

Figure 42: Strain-rate vs maximum applied stress for creep-fatigue data collected at 600C and 650C with a 900s dwell time, creep data from literature is shown for comparison. [32], [100], [101]

Figure 43: Stress-strain hysteresis loops from the primary, secondary and tertiary regimes of a creep-fatigue test with 85 MPa max load and 900s dwell time at 650C.

Figure 44: Creep strain and anelastic strain per cycle as a function of cycle number for the CF curve shown in Figure 39.

Figure 45: Time at maximum load and time at minimum load vs. anelastic strain for specimens tested with the same maximum and minimum dwell loads.

Figure 46: Effect of creep load on anelastic strain and rupture life for CF specimens with matching maximum and minimum load dwell time.

Figure 47: Heterogenous distribution of dislocation content which is characteristic of CSEFS under viscoplastic deformation conditions. (A) BF STEM-DCI showing heterogeneous dislocation content in adjacent subgrains. (B) TKD IPF Map showing low misorientation between subgrains with dissimilar dislocation densities. (C) GROD map illustrating difference in dislocation density (in good agreement with A).

A STEM-ADF B

10 nm

C D

Figure 48: (A) STEM-ADF image showing an edge-on subgrain boundary and holes drilled for EDS drift-correction. (B) EDS map showing Fe deficiency, (C) EDS map showing Cr enrichment, (D) EDS map showing Mo enrichment.

A D B

C

B C D

M23C6 M23C6 M23C6 M23C6 M23C6

Solute Solute Free LAGB Atmosphere

Figure 49: (A) Schematic stress-strain hysteresis showing the subgrain boundary-solute interaction state at different points during the loading cycle. (B) Solute drag occurring during monotonic creep and the later stages of the dwell period during CF deformation. (C) Anelastic backflow allowing boundary dislocations to break free from solute atmosphere. (D) Subgrain migrates without solute atmosphere until boundary dislocations recollect solute.

Figure 50: Strain vs. cycle number for CF tests with different maximum load dwell times, at 650C and 85MPa maximum load.

Figure 51: Anelastic backflow per cycle for specimens with varied maximum load dwell time.

Figure 52: Creep strain per cycle for CF specimens with varied maximum load dwell time.

Figure 53: Average anelastic backflow, creep strain, and net forward strain per cycle as a function of dwell time.

Figure 54: Net strain per cycle vs. cycles to failure for CF specimens with varied maximum load dwell time.

A B

100 µm 400 µm C D

400 µm 5 µm

Figure 55: (A) Fracture surface of a load-controlled CF specimen, showing evidence of ductile fracture and microvoids. (B) Polished fracture surface, showing coalescence of creep voids. (C) Wide field of view (FOV) cross-section of a G91 CF specimen after rupture, showing creep voids and coalescence at the fracture surface. (C) Narrow FOV cross-section, showing fine creep voids at GBs and precipitates.

50m 1m

Figure 56: G91 normalized and tempered microstructure, left: optical micrograph showing PAGB and martensite. Right: STEM-ADF diffraction contrast image showing high dislocation density, fine subgrains and second phase particles at subgrain boundaries.

WM BM

100 µm

CGHAZ FGHAZ ICHAZ

Figure 57: Chemically etched HAZ from a G91 weldment (CMT-as welded)

a b

BM HAZ WM HAZ BM WM B M 1 mm

Figure 58: (a) Chemically etched and hardness indented CMT specimen illustrating method for finding various regions in the HAZ. (b) Microhardness measurements, showing HAZ traverses from the FCAW and CMT weldments.

Constant Load

Loadi ng

Unloa ding

Figure 59: Schematic stress-strain hysteresis loop from a load-controlled CF test on G91.

a b M C 23 6

MX

1 µm 500 nm

Figure 60: STEM-DCI images from the FCAW ICHAZ: (a) ADF image illustrating fine subgrains and M23C6 precipitates at subgrain boundaries. (b) BF image illustrating significant dislocation content within subgrains, M23C6 at subgrain boundaries and MX carbonitrides interacting with mobile dislocations in subgrain boundaries.

FCAW 100 MPa � ̇ = 3.43� 10 s-1 FCAW 85 MPa � ̇ = 1.18� 10 s−1

CMT CMT 100 MPa 85 MPa � ̇ = 2.96� 10 s−1 � ̇ = 8.04� 10 s−1

Figure 61: Creep-fatigue performance comparison for CMT and FCAW weldments.

BM HAZ WM BM HAZ WM

1.1 mm 1.9 mm

FCAW CMT

Figure 62: Optical micrographs showing HAZ width reduction in the CMT specimens when compared to the FCAW specimens.

a HAZ FB WM b HAZ FB WM

FCAW 1 mm CMT 1 mm c d

200 µm 200 µm

Figure 63: Cross-section showing ICHAZ failure in (a) FCAW and (b) CMT weldments. Creep-void coalescence at the fracture surface of (c) CMT and (d) FCAW weldments.

Figure 64: SEM images illustrating precipitate distributions for G91 weldments before and after CF deformation.

Figure 65: Cumulative probability for M23C6 equivalent diameter for weldments before and after CF deformation.

CMT-PWHT CMT-CF 1 µm

FCAW-PWHT FCAW-CF

Figure 66: STEM-ADF images showing low-angle boundary (1-5 misorientation) subgrain structure in CMT and FCAW weldments.

200 nm 200 nm CMT-PWHT CMT-CF

500 nm 500 nm

FCAW-PWHT FCAW-CF Figure 67: Dislocation density after PWHT (Left) and CF-rupture (Right) for CMT weldments (Top) and FCAW weldments (Bottom).

CMT-Sim-PWHT FCAW-Sim-PWHT

500 nm 500 nm

Figure 68: Representative microstructures from simulated ICHAZ specimens, confirming successful ICHAZ microstructure simulations.

Figure 69: CF results from simulated HAZ CF tests performed at 100 MPa and 85 MPa max dwell loads.

Figure 70: Combined full weldment and simulated ICHAZ CF data.

BM ICHAZ WM

After Rupture 0.085 wt%

After PWHT 0.081 wt% t=0s t=23s t=142s t=7342s (PWHT) t=113040s (Failure)

BM ICHAZ WM

After Rupture 0.093 wt%

t=0s After PWHT t=13s 0.086 wt% t=44s t=7255s (PWHT) t=635400s (Failure)

Figure 71: DICTRA simulations showing carbon diffusion from the FZ to the ICHAZ during welding and PWHT of FCAW (Top) and CMT (bottom) weldments.

Figure 72: Distribution of longitudinal strain in a cross-weld specimen of FCAW, which shows a significant strain accumulation in the ICHAZ. Strain map was measured using digital image correlation. Testing temperature = 650 C, peak stress = 100 MPa and dwell time = 900 s.

Figure 73: Comparison of local strain in HAZ for FCAW versus CMT cross-weld specimens tested at 150 MPa. Local strain measured by high-temperature DIC technique.

Figure 74: Distribution of longitudinal strain prior to failure in cross-weld specimens of (a) FCAW and (b) CMT measured by DIC. Testing temperature = 650 °C, peak stress = 150 MPa and dwell time = 900 s.