Heat treatment regime of FGH95 superalloy
\r\n\tThis book aims to explore the issues around the rheology of polymers, with an emphasis on biopolymers as well as the modification of polymers using reactive extrusion.
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He has 20 years experience in waste and by-product valorisation with an emphasis on renewable materials and biological products. Since his tertiary studies, Johan’s knowledge in the engineering field of sustainable products has led to a number of innovative developments in the engineering industry. His research area covers a wide range of topics, such as polymer extrusion, rheology, material properties, protein analysis, chemical modification of proteins as well as protein composites and nano-composites.",institutionString:"University of Auckland",position:null,outsideEditionCount:0,totalCites:0,totalAuthoredChapters:"0",totalChapterViews:"0",totalEditedBooks:"1",institution:{name:"University of Auckland",institutionURL:null,country:{name:"New Zealand"}}}],coeditorOne:{id:"294363",title:"Dr.",name:"Velram",middleName:null,surname:"Mohan",slug:"velram-mohan",fullName:"Velram Mohan",profilePictureURL:"https://mts.intechopen.com/storage/users/294363/images/system/294363.jpeg",biography:null,institutionString:"University of Auckland",position:null,outsideEditionCount:0,totalCites:0,totalAuthoredChapters:"0",totalChapterViews:"0",totalEditedBooks:"0",institution:{name:"University of Auckland",institutionURL:null,country:{name:"New Zealand"}}},coeditorTwo:null,coeditorThree:null,coeditorFour:null,coeditorFive:null,topics:[{id:"6",title:"Biochemistry, Genetics and Molecular Biology",slug:"biochemistry-genetics-and-molecular-biology"}],chapters:null,productType:{id:"1",title:"Edited Volume",chapterContentType:"chapter",authoredCaption:"Edited by"},personalPublishingAssistant:{id:"300344",firstName:"Danijela",lastName:"Pintur",middleName:null,title:"Ms.",imageUrl:"https://mts.intechopen.com/storage/users/300344/images/8496_n.png",email:"danijela.p@intechopen.com",biography:null}},relatedBooks:[{type:"book",id:"1332",title:"Products and Applications of Biopolymers",subtitle:null,isOpenForSubmission:!1,hash:"8dee78e87e2f654541d4285e7cdd5212",slug:"products-and-applications-of-biopolymers",bookSignature:"Casparus Johannes Reinhard Verbeek",coverURL:"https://cdn.intechopen.com/books/images_new/1332.jpg",editedByType:"Edited by",editors:[{id:"102391",title:"Dr.",name:"Casparus",surname:"Verbeek",slug:"casparus-verbeek",fullName:"Casparus Verbeek"}],productType:{id:"1",chapterContentType:"chapter",authoredCaption:"Edited by"}},{type:"book",id:"6694",title:"New Trends in Ion Exchange Studies",subtitle:null,isOpenForSubmission:!1,hash:"3de8c8b090fd8faa7c11ec5b387c486a",slug:"new-trends-in-ion-exchange-studies",bookSignature:"Selcan Karakuş",coverURL:"https://cdn.intechopen.com/books/images_new/6694.jpg",editedByType:"Edited by",editors:[{id:"206110",title:"Dr.",name:"Selcan",surname:"Karakuş",slug:"selcan-karakus",fullName:"Selcan Karakuş"}],productType:{id:"1",chapterContentType:"chapter",authoredCaption:"Edited by"}},{type:"book",id:"1591",title:"Infrared Spectroscopy",subtitle:"Materials Science, Engineering and Technology",isOpenForSubmission:!1,hash:"99b4b7b71a8caeb693ed762b40b017f4",slug:"infrared-spectroscopy-materials-science-engineering-and-technology",bookSignature:"Theophile Theophanides",coverURL:"https://cdn.intechopen.com/books/images_new/1591.jpg",editedByType:"Edited by",editors:[{id:"37194",title:"Dr.",name:"Theophanides",surname:"Theophile",slug:"theophanides-theophile",fullName:"Theophanides Theophile"}],productType:{id:"1",chapterContentType:"chapter",authoredCaption:"Edited by"}},{type:"book",id:"3161",title:"Frontiers in Guided Wave Optics and Optoelectronics",subtitle:null,isOpenForSubmission:!1,hash:"deb44e9c99f82bbce1083abea743146c",slug:"frontiers-in-guided-wave-optics-and-optoelectronics",bookSignature:"Bishnu Pal",coverURL:"https://cdn.intechopen.com/books/images_new/3161.jpg",editedByType:"Edited by",editors:[{id:"4782",title:"Prof.",name:"Bishnu",surname:"Pal",slug:"bishnu-pal",fullName:"Bishnu Pal"}],productType:{id:"1",chapterContentType:"chapter",authoredCaption:"Edited by"}},{type:"book",id:"3092",title:"Anopheles mosquitoes",subtitle:"New insights into malaria vectors",isOpenForSubmission:!1,hash:"c9e622485316d5e296288bf24d2b0d64",slug:"anopheles-mosquitoes-new-insights-into-malaria-vectors",bookSignature:"Sylvie Manguin",coverURL:"https://cdn.intechopen.com/books/images_new/3092.jpg",editedByType:"Edited by",editors:[{id:"50017",title:"Prof.",name:"Sylvie",surname:"Manguin",slug:"sylvie-manguin",fullName:"Sylvie Manguin"}],productType:{id:"1",chapterContentType:"chapter",authoredCaption:"Edited by"}},{type:"book",id:"371",title:"Abiotic Stress in Plants",subtitle:"Mechanisms and Adaptations",isOpenForSubmission:!1,hash:"588466f487e307619849d72389178a74",slug:"abiotic-stress-in-plants-mechanisms-and-adaptations",bookSignature:"Arun Shanker and B. 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And especially, Turbine disk is one of the important component parts for advanced aero-engines, serving usually in the condition of 540~840°C at higher stress, so the turbine disks are required to be of high temperature tolerance and creep resistance (Flageolet B. et al., 2005; Liu D. M. et al., 2006 ). Thereinto, the continuous plastic flow of a material during creep can eventually result in large plastic deformations and significant modifications to the microstructure of the material, so that occurs the creep fracture of the turbine disk parts. Therefore, it is very important for the aero-engine materials to have a better property of the creep resistance.
Traditional wrought superalloys can hardly meet the requirements of the turbine disks in the advanced aerospace for their poor temperature tolerance and loading capacity resulted from their serious composition segregation in ingots and poor hot processability (Park N. K. & Kim I. S., 2001; Lherbier L. W. & Kent W. B., 1990; S. Terzi et al., 2008 ), especially the weaker cohesive force of grain boundaries (Paul L., 1988; Wang P. et al., 2008; Jia CH. CH. et al., 2006 ). While nickel based powder superalloys are an excellent material used for preparing the high temperature rotating section of the advanced aero-engine because its advantages are no macro-segregation in the ingot, chemical composition uniformity and high yield strength (Lu Z. Z. et al., 2005; Zhou J. B. et al., 2002 ).
FGH95 alloy is a nickel-base powder superalloy in which microstructure consists of matrix, and carbide phases. The characteristics of FGH95 superalloy include the high extent of alloying and high volume fraction of -phase (about 50%), besides the alloy possesses excellent integrating mechanical properties at 650°C (Domingue J. A., et al. 1980; Hu B. F. et al., 2006 ). Moreover, various size, morphology and distribution of phase in the alloy can be obtained by different heat treatment regimes (Zainul H. D., 2007 ). The preparation technologies of FGH95 superalloy includes the powder pretreatment, hot isostatic pressing (HIP) and heat treatment. The heat treatment regimes of the alloy include the high temperature solution and twice aging treatment. After solution treated at high temperature, the alloy may adopt the different cooling methods, such as cooled in molten salt or in oil bath, and the microstructure and creep properties of the alloy are related to the chosen heat treatment regimes (Klepser C. A., 1995 ).
The deformation mechanism of the polycrystalline Ni-base superalloys during creep includes twinning, dislocations by-passing or shearing into the phase (Raujol S. et al., 2004 ; Kovarik L. et al., 2009 ). Actually, the mechanical properties and creep behaviors of the alloy are related to the quantities, morphology and distribution of -phase, and especially, the configuration of the boundary and carbides have an important effect on the creep resistance of the alloy (Viswanathan G. B. et al., 2005; Hu B. F. et al., 2003 ). For example, after the alloy is solution treated for cooling in molten salt, the total number and size of secondary phase increase, which can effectively improve the plasticity of the alloy at high temperature ( Zhang Y. W. et al., 2002 ). Because the various microstructures in the alloy may be obtained by different heat treatment regimes, it is very important to understand the influence of heat treatment regimes on the microstructure and creep resistance of the alloy.
In this chapter, the different HIP alloys are solution treated at different temperatures, and cooled in the molten salt or oil bath, respectively, then through twice aging treatment. Besides, some full heat treated alloys are aged for different time at high temperatures, and then the parameters of , phases in the alloy are measured for evaluating the effect of the long term aging time on the misfits. The creep properties of the alloy treated by different heat treated regimes are measured under the conditions of the applied different temperatures and stresses, and the microstructures of the alloy are observed by using SEM and TEM for investigating the influences of the heat treatment regimes on the microstructure and creep properties. Additionally, the deformation mechanism and fracture feature of the alloy during creep are briefly discussed.
FGH95 powder particles of the nickel-base superalloy with the size of about 150 meshes were put into a stainless steel can for pre-treating at 1050 °C for 4 h. The can containing FGH95 powder alloy was hot isostatic pressed (HIP) for 4 h under the applied stress of 120 MPa at 1120 °C, 1150 °C and 1180 °C, respectively. The heat treatment and long term aged treatment regimes of the alloy are listed, respectively, in the Table 1 and Table 2. The cooled rates of the specimen in the oil bath and molten salt are measured to be about 205 °C/min and 110 °C/min, respectively. The error ranges of the used heating furnace are about ± 2 °C. The chemical composition of FGH95 superalloy is shown in Table 3.
HIP Temp., (°C) | Solution Temp., (°C) | Quenching Temp., (°C) | Double aging treatment |
1120 | 1140 for 1 h 1150 for 1 h 1160 for 1 h 1165 for 1 h | cooled for 15 min in molten salt at 583 °C | 870 °C×1h + 650 °C×24h |
1180 | 1150 for 1 h 1160 for 1 h | cooled for 15 min in oil bath at 120 °C | |
1150 for 1 h 1160 for 1 h | cooled for 15 min in molten salt at 583 °C |
Heat treatment regime of FGH95 superalloy
HIP Temp., (°C) | Heat treatment | Aging regime |
1120 | 1155 for 1 h + cooled for 15 min in molten salt at 550 °C + 870 °C×1h +650 °C×24 h | 450 °C×500 h 450 °C×1000 h 550 °C×500 h 550 °C×1000 h |
1150 1180 | ---- |
Long time aging treated regime of FGH95 superalloy
C | B | Cr | Co | Al | Ti | W | Mo | Nb | Ni |
0.060 | 0.012 | 12.98 | 8.00 | 3.48 | 2.55 | 3.40 | 3.40 | 3.50 | Bal |
Composition of FGH95 superalloy(mass fraction,%)
By means of the anode selective dissolving method, the volume fraction of -phase in FGH95 alloy was measured to be about 47%. Thereinto, the electrolytic extraction of phase in the alloy was conducted for separating from the matrix under the condition of the temperature at 0°C and current density about 50mA/cm2. The choosing electrolyte solution consisted of (NH4)2SO4 and citric acid, the experimental device of the electrolytic extraction was shown in Fig. 1. After the electrolytic extraction was conducted, the granularity distribution, phases constituting and the misfit of , phases in the alloy were measured by means of the XRD analysis and SEM/EDS observation.
The ingot of FGH95 superalloy was cut into the specimens with the cross-section of 4.5 mm 2.5 mm and the gauge length of 20 mm, and the size of the sample was shown in Fig. 2. Uniaxial constant load tensile testing was performed, in a GWT504 model creep testing machine, for measuring creep curves under the experimental conditions of 984 MPa ~ 1050 MPa and 630 °C ~ 670 °C. The yield strength of FGH95 alloy was measured to be 1110 MPa at 650 °C. The strain data of the alloy at different conditions were measured with an extensometer to portray the creep curves, twice of the each creep testing were conducted for ensuring the statistical confidence. The specimens of FGH95 alloy at different states were grinded and polished for observing the microstructure by using SEM and TEM, so that the influence of the heat treatment technics on the microstructure, the creep feature and fracture mechanism of the alloy was investigated.
Experimental equipment1--Power supply, 2--sample, 3-- cathode, 4--solution, 5--container
Schematic diagram of the tensile creep sample
After the alloy was hot isostatic pressed at different temperatures, the microstructure of the HIP alloys was shown in Fig. 3. The regions which were encompassed by coarse phase were defined as previous powder particles. After 1120 °C hot isostatic pressing molded, the alloy consisted of and phases, thereinto, the coarser phase distributed around the powder particles was defined as the previous particle boundaries (PPB). Therefore, the configuration of sphere-like previous particle was clearly appeared, and the power particle size was about 15~25 μm, as shown in Fig. 4(a).
With the HIP temperature increased to 1150 °C, the size of the powder particles was similar to the former, but the sphere-like configuration was not clearly. The PPBs consisted mainly of the coarse phase, and the size and amount of the coarser phase decreased slightly as shown in Fig. 4(b). As the HIP temperature increased to 1180 °C, the size of the grain grew up obviously, being about 20~40 μm. Besides, the grain boundaries appeared the straight-like feature, and the amount and size of the coarse phase decreased obviously as shown in Fig. 4(c). The dark regions around the powder particles were defined as the previous particle boundaries (PPB) in which the secondary phase was precipitated along the different orientations as marked by letters A and B.
Microstructure of the alloy after HIP treated at different temperatures. (a) 1120 °C,(b) 1150 °C,(c) 1180 °C
Magnified morphology of the alloy HIP treated at different temperatures. (a) finer -phase precipitated within the grain of the alloy treated by HIP at 1120 °C, (b) after HIP treated at 1150 °C, no - phase particles precipitated in the regions near the coarser -phase as marked by arrow, (c) after the alloy treated by HIP at 1180 °C, particle-like carbides precipitated along the boundary as marked by short arrow.
The magnified morphology of the alloys which were hot isostatic pressed at different temperatures was shown in Fig. 5. Thereinto, the phase displayed the black color due to the dissolved during chemical corrosion, and the matrix which is not dissolved displays the gray color. After HIP treated at 1120 °C, the coarser phase which distributed around the PPB was about 1~2 μm in size and was defined as the primary phase. Besides, the fine secondary phase was regularly distributed along the same orientation within the previous powder particles as shown in Fig. 5(b). As the HIP temperature increased to 1150 °C, the secondary phase about 0.1~0.3 μm in size was dispersedly precipitated within the grain, and the coarser phase still existed in the PPB regions. Moreover, the depleted zone of the fine -phase appeared in the regions near the coarser phase, as marked by the arrow in Fig. 5(b).
The magnified morphology of the powder particle in the 1180 °C HIP alloy was shown in Fig. 5(c), indicating that the fine secondary particles with different orientations were precipitated within the same grain as marked by A and B in Fig. 5(c). No fine particles were precipitated in the PPB regions near the coarser phase, so the region was defined as the depleted zone of the fine -phase as marked by the long arrow in Fig. 5(c). Moreover, some white carbide particles were precipitated in the PPB region as marked with the white short arrow in Fig. 5(c).
The microstructure of the different temperature HIP alloys which were solution treated at 1155 °C, cooled in the molten salt at 520 °C and twice aging treated was shown in Fig. 6. The microstructure of the 1120 °C HIP alloy after full heat treated was shown in Fig. 6(a), illustrating that a few of coarser phase was distributed along the grain boundaries. And the size of the coarse phase was about 1 ~ 2 μm, as marked with the white short arrow in Fig. 6(a). Besides, the fine phase was dispersedly precipitated within the grain as marked with the long arrow in Fig. 6(a). Comparing to Fig. 6(a), the microstructure of the 1150 °C HIP after full heat treatment had no obvious distinction to the former, and the grain size was about 10~25μm. The secondary phase was dispersedly distributed in the alloy as shown in Fig. 6(b). With the HIP temperature increased to 1180 °C and after full heat treatment, the PPB trace was kept in the alloy as marked with the long arrow in Fig. 6(c), indicating that the grain grew up obviously after HIP treated at 1180 °C, and the coarse phase dissolved completely, only kept a few of primary phase which size was about 1μm in the grain boundaries as marked with the short arrow in Fig. 6(c). Moreover, some fine carbide particles were dispersedly precipitated in the alloy as shown in Fig. 6(c).
Microstructure of the different temperatures HIP alloy after fully heat treated. (a) 1120 °C, (b) 1150 °C, (c) 1180 °C
The magnified morphology of the different temperature HIP alloy after fully heat treated was shown Fig. 7. After the 1120 °C and 1150 °C HIP alloys were fully heat treated, a few of coarse phase was precipitated along the boundary regions as marked with short arrow in Fig. 7(a) and (b). And the coarse phase appeared in the boundary as shown in Fig. 7(a).
Magnified morphology of the different temperature HIP alloy after fully heat treated (a) 1120 °C, (b) 1150 °C, (c) 1180 °C
Moreover, fine white particles were discontinuously precipitated along the boundaries as marked with the long arrow in Fig. 7(a) and (b), and the fine secondary phase and white particles were dispersedly distributed in the alloy as shown in Fig. 7(a) and (b). With the HIP temperature increased to 1180 °C and after fully heat treated, the sphere-like PPB kept in the alloy as marked with long arrow in Fig. 7(c), and the boundaries displayed an obvious straight feature as marked with short arrow in Fig. 7(c). Moreover, fine phase was dispersedly distributed within the grain, and the white particles of about 0.2 μm in size were dispersedly precipitated in the grain as shown in Fig. 7(c).
When the FGH95 alloy was HIP treated at 1120 °C and 1150 °C which is lower than the dissolving temperature of phase ( Tm= 1160 °C ), the coarse phase distributed around the PPB were not completely dissolved. Therefore, the phase which was reserved in the PPB regions was going to grow up during the slowly cooling stage of HIP, which resulted in forming the morphology of the coarser particles distributed along the PPB. The size of previous powder particles were kept in the alloy because no enough deformed energy which come from the HIP deformation promoted the powder particles growth up, as shown in Fig. 4(a) and (b).
When the alloy was hot isostatic pressed under the conditions of 120 MPa at 1180 °C, the coarser phase in the PPB regions was completely dissolved, which may enhance the diffusion rate of the elements and reduce the transferring resistance of the boundaries to promote the growth of the grain, therefore, the alloy displayed a larger grain size as shown in Fig. 4(c). At the same time, the nucleation and growth of the carbide occurred on the fine oxides which existed in the powder particle surface, and the carbide nucleation process only needs lower energy (Domingue J. A. et al., 1980). Consequently, some white carbide and coarser particles were discontinuously precipitated in the PPB regions as shown in Fig. 5(c), and the depleted zones of the fine phase appeared in the region near the coarse particles, due to the coarsening of particles to consume the much more elements of former, as shown in Fig. 5(c).
After the 1120 °C HIP alloy was solution treated at 1140 °C and twice aged, the grain size of the alloy was homogeneous, and fine phase was dispersedly precipitated within the grain, as shown in Fig. 8(a). Besides, much more coarse phase precipitated along the grain boundaries as marked with the arrow in Fig. 8(a). The magnified morphology was shown in Fig. 8(b), significant amount of fine phase were dispersedly distributed in the alloy, and the depleted zone of -phase appeared around the coarse phase which distributed along the boundary regions, as shown in the Fig. 8(b).
With the solution temperature increased to 1150 °C, the microstructure of the alloy consisted of the and phases, the average grain size of the alloy was about 10~20 μm, as shown in Fig. 9(a), which indicated that some coarser -particles are precipitated in the wider boundary regions, and the average size of the coarser phase was about 1~2.5 μm. The magnified morphology of the alloy was shown in Fig. 9(b), indicating that significant amount of the fine particles were dispersedly distributed within the grains, the size of the
Microstructure of the alloy after solution treated at 1140 °C. (a) after solution treated at 1140 °C, coarse phase distributed in the alloy, (b) no finer phase precipitated in the regions near the coarser phase
ones was about 0.1~0.3 μm. And no fine -particles were precipitated in the regions near the coarser -phase, the regions were defined as the depleted zone of the fine -phase as marked by the arrow in Fig. 9(b), the magnified morphology of the depleted zone of -phase was marked by the arrow in Fig. 9(c).
Microstructure of the alloy after solution treated at 1150 °C. (a) After solution treated at 1150 °C, some coarser phase precipitated in the wider boundary regions, (b) fine phase distributed dispersedly within the grains, and the depleted zones of the fine phase marked by arrow, (c) magnified morphology of the depleted zone of -phase
As the solution temperature increased to 1160 °C, the grain boundaries appeared obviously, and the average size of the grains in the alloy was about 15 ~ 25 μm as shown in Fig. 10(a). Compared to Fig. 9(a), the coarser -precipitates in the boundary regions disappeared, and the white particles with size about 0.2 μm were precipitated within the grains and along the boundaries as marked by the arrows in Fig. 10(a). The magnified morphology of the alloy was shown in Fig. 10(b), indicating that the coarser particles and depleted zone of the fine phase had disappeared, and the secondary phase was dispersedly distributed within the grains. The grain boundary in the alloy was marked by longer arrow and some particles were homogeneously precipitated along the boundaries and within the grains as marked by the short arrows in Fig. 10(b).
Microstructure of the alloy after solution treated at 1160 °C. (a) after solution treated at 1160 °C, the coarser phase disappearing and the grain boundaries appearing obviously in the alloy, (b) the secondary phase distributed dispersedly within the grain, and some particles precipitated in the alloy as marked by the arrows
Microstructure of the alloy after solution treated at 1165 °C. (a) After solution treated at 1165 °C, the average size of the grains in the alloy increasing obviously, and films of white phase precipitated along the boundaries as marked by arrows, (b) the secondary phase distributed dispersedly within the grains, and films of the white phase distributed along the grain boundaries
As the solution temperature increased to 1165 °C, the grain sizes in the alloy increased, the linear-like boundaries appeared in the alloy, and the films of the white phase were continuously precipitated along the boundaries as marked by the arrows in Fig. 11(a). The magnified morphology of the white phase was shown in Fig. 11(b), indicating that white particles were continuous precipitated to form the films along the boundaries as marked by the arrow in Fig. 11(b), and significant amount of the fine secondary phase were precipitated within the grains, no depleted zones of the fine phase were detected in the alloy. The grain sizes after the alloy was solution treated at various temperatures were measured as listed in Table 4. Moreover, by means of composition analysis under SEM/EDS, it is indicated that the elements Nb, Ti and C are richer in the white particles which were located within the grain and boundary regions as shown in Fig. 10 and Fig. 11, respectively.
After the alloy was solution treated at 1160 °C, the fine -particles with the size of about 0.1 μm were dispersedly precipitated within the grains as shown in Fig. 12(a), the particles can effectively hinder the dislocation movement to enhance the creep resistance of the alloy. The smaller space between the fine -particle was measured to be about 0.03 μm, the bigger space between the fine -particles is measured to be about 0.12 μm as marked by letters L1 and L2 in Fig. 12(a), respectively. It was indicated by TEM/EDS analysis that the elements Nb, Ti and C richer in the carbide particle which was located in the boundary as shown in Fig. 12(b), and the particle was identified as (Nb, Ti)C phase by means of the diffraction spots analysis as marked in Fig. 12(c).
Solution temp. (°C) | 1140 | 1150 | 1160 | 1165 |
Average grain sizes (μm) | 10~20 | 10~20 | 15~25 | 20~40 |
Grain sizes of the alloy solution treated at different temperatures
Morphologies of and carbide phases. (a) Fine phase precipitated dispersedly within the grain, (b) carbide particles precipitated along the boundaries, (c) SAD patterns
After the 1180 °C HIP alloy was solution treated at different temperatures and cooled in oil bath at 120 °C and twice aging treated, the alloy displayed the various sizes of grains in which the fine phase was dispersedly and regularly precipitated, as marked by white long arrow in Fig. 13(a). The bigger grains was about 40~60 μm in size, the smaller grains was about 20 μm in size, the bunch-like coarser particles were distributed in the PPB regions, or some coarser particles were congregated in the local regions as marked by black arrow in Fig. 13(a). The depleted zone of the fine -phase appeared in the regions near the coarser phase as marked with the white short arrow in Fig. 13(a). As the solution temperature increased to 1160 °C, the coarser phase in the alloy was completely dissolved due to the higher solution temperature. Moreover, the fine phase and the white particles were dispersedly distributed within the grains, respectively, and some white particles were precipitated along the boundaries as marked by white arrow in Fig. 13(b). Furthermore, the twinning appeared within the grain as marked by black arrow in Fig. 13(b).
Morphology of the alloy after solution treated at different temperatures and cooled in oil bath. (a) After solution treated at 1150 °C, the fine phase distributed dispersedly within the grains, and the depleted zone of -phase as marked by shorter arrow. (b) after solution treated at 1160 °C, the fine phase and carbide particles distributed dispersedly within the grains, respectively
After solution treated at 1150 °C, then cooled in molten salt and twice aged, the microstructure of the alloy consisted of the particle-like phase with different size, the coarser -phase was distributed in the PPB regions as marked by white arrow, and the fine phase was dispersedly distributed within the grain as marked by black arrow in Fig. 14(a). Furthermore, the depleted zone of the fine phase appeared in the region near the coarser -phase as marked by white short arrow in Fig. 14(a), and the size of the fine phase was about 0.15 μm.
Morphology of the alloy after solution treated at different temperatures, and cooled in molten salt. (a) After solution treated at 1150 °C, the fine phase distributed dispersedly in the grains, and the coarse particles precipitated along the boundaries, (b) after solution treated at 1160 °C, straight-like boundaries marked by black arrow, and the fine particles distributed dispersedly within the grains.
After the 1120 °C HIP alloy was solution treated at 1155 °C, cooled in molten salt at 520 °C and twice aging treated, the microstructure of the full treated alloys which were kept at 450 °C and 550 °C for 500 h respectively was shown in Fig. 15(a) and (b). Compared to the alloy which was not long time aged, no obvious distinction on the microstructure was detected in the alloy which was long time aged for 500 h at 450 °C, as shown in Fig. 15(a). Thereinto, no obvious change on the grain size was found, and a few coarse particles about 1~2 μm in size was precipitated along the boundaries as shown in the Fig. 15(a). With the long time aging temperature increased to 550 °C, the grain size about 15~20 μm was similar to the former, and some coarse phase were distributed along the boundaries as shown in Fig. 15(b).
Microstructure of the alloy aged for 500 h at different temperatures. (a) 450 °C, (b) 550 °C
Magnified morphology of the alloy aged for 500 h at different temperatures. (a) 450 °C, (b) 550 °C
The magnified morphology of the alloys aged for 500 h at 450 °C and 550 °C was shown in Fig. 16(a) and (b), respectively. After the alloy aged for 500 h at 450 °C, fine white carbide particles were precipitated along the boundaries as marked with short arrow in Fig. 16(a). Moreover, the secondary phase about 0.1~0.2 μm in size was dispersedly distributed in the grain as marked with the long arrow in Fig. 16(a). With the temperature increased to 550 °C, the microstructure of the alloy was similar to the former, and the fine carbide particles about 0.1~0.25 μm in size were distributed along the boundaries as marked with the short arrow in Fig. 16 (b). Besides, the fine precipitates were dispersedly distributed in the alloy as marked with the long arrow in Fig. 16(b).
The microstructure of the full heat treated alloys aged for 1000 h at 450 °C and 550 °C was shown in Fig. 17(a) and (b), respectively. Fig. 17(a) showed that the microstructure was similar to the one of the aged-free alloy, and the grain size was about 15~20 μm. Moreover, a few coarse phase was precipitated along the grain boundaries, and the size of precipitates was about 1~2 μm as shown in Fig. 17(a). As the aging temperature increased to 550 °C and kept for 1000 h, no obvious distinction on the size of the grain and coarse phase was detected, as shown in Fig. 17(b).
Microstructure of the alloy aged for 1000 h at different temperatures. (a) 450 °C, (b) 550 °C
Magnified morphology of the alloy aged for 1000 h at different temperatures. (a) 450 °C, (b) 550 °C
After aged for 1000 h at 450 °C and 550 °C, the magnified morphology of the alloy was shown in Fig. 18. Figure 18 shown that, after aged for 1000 h at 450 °C, some fine white carbide particles were precipitated along the boundaries as marked by the short arrow in Fig. 18 (a). Besides, twinning appeared obviously in the alloy as marked with the long arrow in Fig. 18(a). With the aging temperature increased to 550 °C, the microstructure of the alloy was similar to the former, and some fine white carbides distributed along the boundaries as marked with the short arrow in Fig. 18(b). Moreover, the secondary phase about 0.1~0.2 μm in size was dispersedly precipitated in the grain as marked with the long arrow in Fig. 18(b).
After FGH95 alloy was full heat treated, the microstructure of the alloy consisted of -matrix, phase and carbide particles. Thereinto, the ordered phase was coherently embedded in the matrix. Besides, both and phases have the FCC structure and close lattice parameters. Moreover, the coherent interfaces was kept in between the / phases, and the certain misfits occurred in the interfaces of / phases. The equations of calculating the parameters and misfit may be expressed as follow:
Where, d is crystal plane distance,\n\t\t\t\t\t is diffraction angle, ( = 0.154 nm, α is lattice parameter, (h, k, l) is crystal plane indexes and is the lattice misfit.
Dependence of aging time and lattice parameters of / phases at different temperatures. (a) 450 °C, (b) 550 °C
After the fully heat treated alloys aged for 500 h, 750 h, 1000 h and 1500 h at 450 °C and 550 °C, the XRD curves of the alloys were measured, in which the located angles were chosen to separate the diffraction peaks of and phases for measuring the crystal plane distance of and phase in the alloy. In the further, the lattice parameters and misfits of the alloys were calculated according to the equations (Eq. 1) and (Eq. 2), and the dependence of the lattice parameters and aged time at different temperatures was shown in Fig. 19. Figure 19(a) shown that the parameters of and phases increased gradually with the aging time at 450 °C, and the parameter increment of phase was more than the one of phase. With the aging temperature increased to 550 °C, the trend of the parameter increment was similar to the former, but the parameter increment of and phases was larger than the former, as shown in Fig. 19(b).
Relationship between the lattice misfit and aging time at different temperatures
The relationship between the lattice misfit and aging time at different temperatures was shown in Fig. 20. This indicated that the lattice misfits decreased from 0.3734% to 0.2030% as the aging time prolonging at 450 °C as marked with letter A in Fig. 20. With the aging temperature increased to 550 °C, the lattice misfit decreased from 0.3734% to 0.2002% as marked with letter B in Fig. 20. Thereinto, the bigger difference between the lattice misfits of the alloys aged at 450 °C and 550 °C appeared in the periods of aging for 500 h and 1000 h.
After the 1120 °C HIP alloy solution treated at 1160 °C and cooled in molten salt at 583 °C and twice aging treated, creep curves of the alloy under different conditions are shown in Fig. 21. Fig. 21(a) was the creep curves measured under different applied stresses at 650 °C. Under the applied stress of 1020 MPa, the alloy displays the shorter initial creep stage and longer steady creep stage, besides, the alloy displays the lower strain rate and longer creep lifetime about 156 h. With the applied stress enhanced to 1034 MPa, the strain rate of the alloy during steady state creep increased, and the lifetime of the alloy decreased to 104 h. As the applied stress enhanced to 1050 MPa, the creep lifetime of the alloy decreased sharply to 51 h. This indicates that, in the ranges of the applied stresses, the alloy is not sensitive to the applied stress at 650 °C.
After fully heat treated, creep curves of 1120 °C HIP alloy at different conditions(a) Applied various stresses at 650 °C, (b) 1034 MPa at various temperatures
The creep curves of the alloy under the conditions of 1034 MPa and different temperatures are shown in Fig. 21(b), illustrating that the alloy possesses the lower strain rate and longer creep lifetime under the applied stress of 1034 MPa at 650 °C, and the strain rate of the alloy during steady state creep was measured to be about 0.00367%/h, the creep lifetime of the alloy was about 104 h. With the creep temperature increased to 660 °C, the strain rate of the alloy during steady state creep increased to 0.00825%/h, and the creep lifetime decreased sharply to 40 h. As the temperature increased to 670 °C, the strain rate during steady state creep further enhanced to 0.01792%/h, and the lifetime of the alloy decreased to 17 h. This indicates that the alloy has an obvious sensitive to the applied temperature.
Under the applied stress of 1034 MPa at 650 °C, creep curves of the alloy solution treated at different temperatures were shown in Fig. 22. The creep curve of the alloy solution treated at 1150 °C was marked by letter A in Fig. 22, through which the strain rate of the alloy during steady state creep was measured to be 0.0102%/h, the lasting time was about 40 h, and the creep lifetime of the alloy was measured to be 67 h. The creep curve of the alloy solution treated at 1160 °C was marked by letter B in Fig. 22, indicating that the alloy displays a lower strain rate during steady-state creep, and the creep lifetime of the alloy was measured to be 104 h. When the solution temperature increased to 1165 °C, the creep lifetime of the alloy was measured to be about 9 h, as marked by letter C in Fig. 22.
After solution treated at different temperatures, the creep lifetimes of the alloy under the applied various stresses and temperatures are measured as listed in Table 5. It may be understood from Table 5 that, under the applied stress of 1034 MPa at 650 °C, the lifetime of the alloy solution treated at 1150 °C was measured to be 67 h. As the solution temperature enhanced to 1160 °C, the lifetime of the alloy increased to 104 h. When the solution temperature enhanced further to 1165 °C, the lifetime of the alloy decreased rapidly to 9 h, as shown in Table 5. This indicates that the solution temperature has an obvious influence on the creep lifetimes of the alloy.
Under the applied stress of 1034 MPa at 650 °C, creep curves of the alloy solution treated at different temperatures
After solution treated at 1150 °C, the lifetime of the alloy under the applied stress of 1020 MPa at 650 °C was about 95 h, and the strain rate during steady state creep was measured to be about 0.00827%/h. The lifetime of the alloy decreased to 67 h as the applied stress increased to 1034 MPa at 650 °C, and the creep lifetime of the alloy decreased to 37 h as the applied stress increased to 1050 MPa at 650 °C. And the alloy solution treated at 1160 °C displays also the similar regularity under the experimental conditions. This indicates that the alloy has an obvious sensitive to the applied stresses and temperatures.
Effect of solution temperature on stress rupture properties of FGH95 alloy
After the alloy was solution treated at 1150 °C and cooled in oil bath at 120 °C, the creep curves of the alloy at different conditions were measured as shown in Fig. 23. Under the applied stress of 1034 MPa at different temperatures, the creep curves of the alloy were shown in Fig. 23(a), it indicated that the alloy possessed the lower strain rate and longer creep lifetime of about 380 h at 630 °C.
Creep curves of the alloy cooled in oil bath at different conditions. (a) Creep curves of the alloy under the applied stress of 1034 MPa at various temperatures, (b) creep curves of the alloy under the applied different stresses at 650 °C
As the creep temperature increased to 640 °C, the lifetime of the alloy decreased to 144 h. When the creep temperature increased to 650 °C, the lifetime of the alloy decreased rapidly to 32 h. This indicates that the alloy had an obvious sensitive to the applied temperatures. While the creep curves of the alloy under the applied different stresses at 650 °C were measured as shown in Fig. 23(b). Under the applied stress of 984 MPa, the alloy displayed a shorter initial creep stage and a longer steady state stage of creep, and the creep lifetime of the alloy was measured to be about 260 h. The lifetime of the alloy decreased to 205 h as the applied stress enhanced to 1010 MPa, but the creep lifetime of the alloy at 650°C/1034 MPa deceased rapidly to 32 h. This indicates that the alloy possessed an obvious sensitive to the applied stress when the applied stress was over 1010 MPa.
Creep curves of the alloy cooled in molten salt at 583 °C at different conditions. (a) Creep curves of alloy under applied stress of 1034 MPa at various temperatures, (b) creep curves of alloy under the applied different stresses at 650 °C
After the solution treated at 1160 °C and cooled in molten salt at 583 °C, the creep curves of the alloy at different conditions were shown in Fig. 24. Under the applied stress of 1034 MPa at different temperatures, the creep curves of the alloy were measured as shown in the Fig. 24(a), indicating that the alloy displayed a lower strain rate and creep lifetime of about 145 h at 650 °C. As the temperature enhanced to 660 °C and 670 °C, the strain rates of the alloy during the steady state creep increased, and the creep lifetimes decreased to 60 h and 32 h, respectively. This indicates that the alloy has an obvious sensitive to the applied temperatures.
Under the applied stress of 1034 MPa at 650 °C, creep curves of the alloy treated by different regimes
The creep curves of the alloy under the applied different stresses at 650 °C were shown in Fig. 24(b). The creep feature of the alloy under the applied stress of 1020 MPa displayed a shorter initial stage and longer steady state stage, and the creep lifetime of the alloy was measured to be 213 h. As the applied stress enhanced to 1034 MPa and 1050 MPa, the creep lifetimes of the alloy were measured to be 145 h and 68 h, respectively. This displays only a smaller decreasing extent of the creep lifetimes, therefore, it may be concluded that no obvious sensitivity is displayed in the alloy in the ranges of the applied stresses.
Creep properties of FGH95 alloy treated by different regimes at 1034 MPa /650 °C
Under the applied stress of 1034 MPa at 650 °C, the creep curves of the alloy cooled in different mediums were shown in Fig. 25, indicating that, compared to the oil cooled alloy, the alloy cooled in molten salt displayed a lower stain rate and longer creep lifetime. Moreover, the creep properties of the alloy treated by different regimes under different conditions are measured as listed in Table. 6, this indicates that the alloy solution treated at high temperature possesses the better creep properties.
Transient strain of the alloy occurs when the loading is applied at high temperature, and the strain rate of the alloy decreases as the creep goes on. The strain rate keeps constant once the creep enters the steady state stage, therefore, the strain rate of the alloy during steady state creep may be expressed by Dorn’s law as follows:
where
After the 1120 °C HIP alloy solution treated at various temperatures, the creep curves of the alloy were measured in the ranges of 650 °C ~ 670 °C and 1020 MPa ~ 1050 MPa. The dependence of the strain rates of the alloy during steady-state creep on the applied temperatures and stresses was shown in Fig. 26. Figure 26(a) showed the relationship between the strain rates and the temperatures under the applied stress of 1034 MPa, and the dependence of the strain rates on the applied stresses at 650 °C was shown in Fig. 26(b). According to the data during the steady state creep, the creep activation energies and stress exponents of the alloys, which were solution treated at 1150 °C and 1160 °C, were calculated to be QA= 510.1 20 kJ/mol, QB = 580.3 20 kJ/mol and nA=15.4, nB=14.1, respectively.
Dependence of the strain rates during steady state creep on the applied temperatures and stresses for the alloy solution treated by different temperatures. (a) Strain rate vs. temperatures at 1034 MPa, (b) strain rate vs. the applied stress at 650 °C
After the 1180 °C HIP alloy was solution treated and cooled in different mediums, the dependences of the strain rates during steady state creep on the applied temperatures and stresses in the ranges of 650 °C ~ 670 °C and 1020 MPa ~ 1050 MPa are shown in Fig. 27. The relationship between the strain rates and the temperatures under applied stress of 1034 MPa was shown in Fig. 27(a), according to the data, the creep activation energies of the alloy treated by different regimes were calculated to be QA= 381.1 20 kJ/mol and QB = 590.3 20 kJ/mol, respectively. The dependence of the strain rates on the applied stresses at 650 °C was shown in Fig. 27(b), according to the data, the stress exponents of the alloy treated by different regimes were calculated to be nA = 17.9 and nB = 13.8, respectively. The values of the creep activation energies and stress exponents were measured as listed in Table 7. It may be concluded according to the data in Table 7 that the alloy cooled in molten salt displayed a better creep resistance.
Relationship between the strain rates and the applied temperatures, stresses for the alloy treated by different regimes. (a) Strain rates & temperatures, (b) strain rates & the applied stresses
Activation energies and stress exponents during the steady-state creep of the alloy treated by different regimes
After the 1120 °C HIP alloy solution was treated at 1160 °C, the morphology of the alloy crept for 104 h up to rupture under the applied stress of 1034 MPa at 650 °C was shown in Fig. 28. Figure 28(a) displayed the feature of the stacking fault as marked by white arrow, and two (1/3)<112> super-Shockleys partials were located on the two sides of the stacking fault. Some dislocation loops with various sizes appeared clearly in the matrix of the alloy, the smaller dislocation loop was marked with black arrow in Fig. 28(a), and larger dislocation loop was located in the region above the stacking fault. It could be deduced by analysis that the <110> super-dislocation shearing into the or -phase may be decomposed to form the configuration of (1/6)<112> shockleys dislocations or (1/3)<112> super-Shockleys partials plus the stacking fault.
Some blocky carbide particles were precipitated in the local area of the alloy as marked by black arrow in Fig. 28(b). The grain boundary in the alloy was marked by white arrow in Fig. 28(c), and dislocations tangles were piled up in the region B near the boundary, which suggested that the boundaries may effectively hinder the dislocation movement during creep. Besides, the morphology of <110> super-dislocation shearing into the -phase was marked by the black arrow, and the stacking fault formed from the dislocation decomposition was marked by the letter A in Fig. 28(c). The blocky carbides were precipitated along the boundary as marked by the black arrow in Fig. 28(d), and some dislocations were piled up in the regions near the carbides, which indicated that the carbide particles can effectively hinder the dislocation movement. But, the facts that significant amount of dislocations were piled up in the regions near the boundaries can cause the stress concentration to promote the initiation and propagation of the micro-crack along the boundaries as the creep goes on.
Microstructure of the alloy after (solution treated at 1160 °C) crept for 104 h up to fracture at 650 °C/1034 MPa. (a) dislocation loops in the alloy, (b) carbide particles precipitated within the grain, (c) morphology of the super-dislocation shearing into phase as marked by black arrow, the stacking fault in the region A and dislocation tangles piled up in the B region near the boundary, (d) carbide particles precipitated along the boundary, and dislocation tangles piled up in the regions near the boundaries
After the 1180 °C HIP alloy solution treated at 1150 °C and cooled in the oil bath, the microstructures of the alloy crept for 260 h up to fracture under the applied stress of 984 MPa at 650 °C was shown in Fig. 29. The fine sphere-like particles about 0.1~0.2 µm in size were dispersedly distributed in the alloy as shown in Fig. 29(a), which may hinder the dislocation movement for improving the creep resistance of the alloy. During creep, the double orientation slipping of dislocations were activated in the local regions of matrix as shown in Fig. 29(b), the deformed dislocations in the form of the tangles were arranged along the level direction, the bundle-like dislocations were arranged along the upright-like direction as marked by arrows, respectively, in Fig. 29(b). During creep, the slipping dislocations moving to the boundary were stopped as marked by white arrow in Fig. 29(c), indicating that the boundary had an obvious effect on hindering the dislocations movement. Therefore, it may be concluded that the deformation feature of the alloy during creep is the double orientations slipping of dislocations activated in the matrix phase.
After solution treated at 1150 °C and cooled in oil bath, microstructure of the alloy crept for 260 h up to fracture under the applied stress of 980 MPa at 650 °C. (a) Fine phase precipitated dispersedly within the grains, (b) double orientations slipping of dislocation as marked by arrows, (c) dislocations slipping stopped on the boundary
After the 1180 °C HIP alloy was solution treated at 1160 °C and cooled in the molten salt, the microstructures of the alloy crept for 72 h up to fracture under the applied stress of 1034 MPa at 650 °C was shown in Fig. 30. The carbide particles were precipitated along the boundaries as marked by arrow in Fig. 30(a), which resulted in the boundary with the uneven feature. The dislocation-free region appeared in the left side of the boundary, significant amount of dislocations were congregated in the another side of the boundaries, and the carbide particle was precipitated along the boundary as marked by the arrow in Fig. 30(b). It may be analyzed according to Fig. 30(b) that the boundaries and the carbides precipitated along the boundaries may hinder the dislocations movement for improving the creep resistance of the alloy, so that significant amount of dislocations are piled up in the right-side region of the boundary as shown in Fig. 30(b).
After solution treated at 1160 °C and cooled in molten salt, microstructure of the alloy crept up to fracture at 1034 MPa / 650 °C. (a) Carbide particles precipitated along the boundary, (b) dislocation tangles piled up in the regions near the boundaries, (c) stacking fault and dislocations appeared within the grain
In the local region of the alloy, the double orientation slipping of dislocations were activated within the grain, the boundary was marked by arrow in the bottom of Fig. 30(c), the directions of the double orientation slipping of dislocations were marked with arrows in Fig. 30(c). The arrow which points to left side corresponds to the orientation of the dislocations tangle, the arrow which points to right-side corresponds to the direction of the stacking fault stripes. It may be deduced by analysis that the stacking fault is formed in between two (1/3)<112> super-Shockleys partials which originates from the decomposition of <110> super-dislocation. The configuration of the partial dislocation + stacking fault can restrain the cross-slipping of the dislocations to improve the creep resistance of the alloy.
Twinning in the alloy during creep at 650 °C/1034 MPa. (a) twinning, (b) SAD patterns
After the alloy was crept up to fracture under the applied stress of 1034 MPa at 650 °C, the feature of the twinning deformation was shown in Fig. 31(a), and the twinning plane was identified as (111) plane by means of the diffraction pattern analysis as shown in Fig. 31(b). The blocky carbide was precipitated along the grain boundary as marked by shorter arrow in Fig. 32(a), and some micro-twinning was detected in the alloy as marked by black arrow. In another local area, the morphology and diffraction spots of the micro-twinning were shown in Fig. 32(b) and (c), this is well agreed with the result in literature (Unocic R. R. et al., 2008 ). The one end of the micro-twinning was stopped at the grain boundary, which suggests that the grain boundaries had an obvious effect on hindering the twinning deformation.
Morphologies of the carbide precipitated along the boundary and micro-twinning formed in the alloy during creep under the applied stress of 1034 MPa at 650 °C. (a) Carbide precipitated along the boundary, (b) micro-twinning, (c) SAD patterns
Significant amount of the fine particles are precipitated within the grains, which may effectively hinder the dislocation movement. When the deformed dislocations move over the phase during creep, the dislocation loops are kept around the particles as shown in Fig. 28(a), which suggests that the deformation feature of the alloy during creep is the dislocations moving over the -phase by Orowan bypassing mechanism. It is reasonable consideration that the various spaces between the particles appear in different regions, the dislocations may bow out along the wider channels between the two particles during creep, and the bowing dislocations move over the -phase by Orowan mechanism to encounter for forming the dislocation loops, as marked with black arrow in Fig. 28(a). When the dislocations bow out along the channels during creep, the applied stress which is enough to overcome the Orowan resistance can be expressed as follows:
Where, is the shear modulus, b is the Burgers vector, and L is the space between two particles. This indicates that the resistance of the dislocation movement increases with the diminishing of the space between the -particles, and the resistance of alloy enhances with the volume fraction of -particles.
Schematic diagram of dislocation bowing out in the congregated regions of the particles
The space between the particles diminishes when some particles are congregated together, the smaller space is measured to be L1 = 0.03 μm as shown in Fig. 12(a), but the bigger space between the -particles is measured to be L2 = 0.12 μm. Compared to the bigger space, the dislocations bowing out in the regions with smaller space needs fourfold Orowan resistance according to the formula (Eq. 6), therefore, it is difficult for the dislocations to bow out along the channels with smaller space, but they can bow out along the channels with larger space to form the larger loops as shown in Fig. 28(a). When the dislocations bow out along the channels with bigger space (L), the formation process of the bigger dislocation loops is schematically shown in Fig. 33.
In the later stage of creep, significant amount of dislocations move to the region near the particles to generate stress concentration, and when the stress value originated from the stress concentration exceeds the yield strength of phase, the <110> super-dislocation may shear into the -phase, as marked by black arrow in Fig. 28(c). Furthermore, the <110> super-dislocations may be decomposed to form the configuration of two (1/3)<112> super-Shockleys partials and the stacking fault, as marked by letter A in Fig. 28(c). The critical stress of dislocation shearing into phase increases with the yield strength. And the critical stress (τcs) can be expressed as follows (Zhang J. SH., 2007) :
Where T is the dislocation line tension, r is the radius of the -particle, b is the Burgers vector, f is the volume fraction of phase, and APB is the antiphase boundary energy originated from the dislocation shearing phase. It may be understood from equation (6) that the critical stress of dislocation shearing into phase increases with the size (r), volume fraction (f) of -phase and antiphase boundary energy (APB) to improve the creep resistance of the alloy.
The observed shearing process spawned the model (Unocic R. R. et al., 2008 ) is illustrated in Fig. 34. The process of shearing the larger particles is presumed to occur by the cooperative movement of the coupled (a/2)[112] dislocation (3D), each of which is dissociated into three (a/3)[112] partials (D).
Model of precipitate shearing by coupled Shockley partials for creating SISF/SESF pairs. After Hirth and Lothe (Unocic R. R. et al., 2008, as cited in Hirth J. P. & Lothe J., 1968 )
The a<112> dislocations are hypothesized to originate from the interaction of two different a<101> super- dislocations originating from different slip systems. For example:
Clearly, this model then requires a high symmetry orientation such that two slip systems experience a relatively large shear stress.
In situ deformation at higher temperature gives rise to a distinctly different mode of shearing in which the extended faults propagate continuously and viscously through both particles and matrix. These extended faults are associated with partials that move in a correlated manner as pairs. Koble (Koble M., 2001 ) induced that these partials may be a/6<112> partials of the same Burgers vector, and that they may be traveling in parallel {111} planes, as illustrated in Fig. 35. Without detailed confirmation of this hypothesis, Kolbe further deduced that these were in fact micro-twins, and that the temperature dependence of the process may be associated with recording that would ensure in the wake of twinning a/6<112> partials as they traverse the particles. The shear strain rate can be expressed as follow:
Where, гpt is the energy of two layered pseudo-twin, and bpt is the magnitude of the Burgers vector of the twinning partials, гtt is the energy of two layered true twin, pt is the density of mobile twinning partials, Dord is the diffusion coefficient for ordering, x is the short range diffusion length (assumed to be several nearest neighbor distances, or ~2b), f2 is the volume fraction of the secondary precipitates, f3 is the volume fraction of the tertiary precipitates. And the effective stress (eff), in the presence of tertiary precipitates, is given by:
The experimental values of parameters such as dislocation density pt, volume fraction of the secondary precipitate that are critical to the prediction can be determined directly from TEM observations. Disk alloys in this temperature regime typically exhibit the creep curves having a minimum rate, with a prolonged increase of creep rate with time. As the fine phase volume fraction decreases during thermal exposure, it is possible that the operation of 1/2[110] matrix dislocations becomes increasingly important. The coarse microstructure (small value of f3) resulting from a slow cooling rate, the deformation is dominated by 1/2<110> dislocation activated in the matrix, and SESF shearing in the secondary precipitates.
Schematic representation of micro-twinning mechanism from shear by identical Shockley partials (D) transcending both the matrix and precipitate in adjacent {111} planes which then require atomic reordering in to convert stacks of CSF into a true twinned structure. After Kolbe (Koble M., 2001 )
After the 1120 °C HIP alloy was solution treated at 1150 °C and isothermal quenched in molten salt at 583 °C, the morphology of the alloy crept for different time under the applied stress of 1034 MPa at 650 °C was shown in Fig. 36. The applied stress direction was marked with the arrow in Fig. 36(a), after the alloy was crept for 40 h, some slipping traces appeared on the surface of the sample, and some parallel slipping traces were displayed within the same grain. Moreover, the various orientations of the slipping trace appeared within the different grains. Besides, the kinking of the slipping traces appeared in the region of the boundaries as marked by arrow in Fig. 36(a). After crept for 67 h up to rupture, the surface morphology of the alloy was shown in Fig. 36(b), indicating that the amount of the slipping trace increased as the creep went on, and the slipping traces were deepened to form the slipping steps on the surface of the specimen. Moreover, the bended slipping traces appeared in the boundary regions, as marked by longer arrow in Fig. 36(b), which was attributed to the effect of the flow metal in the -free phase zone where is lower in strength. Besides, the cracks were initiated in the distortion regions of slipping traces as marked by shorter arrow in Fig. 36(b).
Surface morphology of the alloy crept for different time up to fracture. (a) After crept for 40 h, a few slipping traces appeared within the different grains, (b) after crept up to fracture, significant amount of the slipping traces appeared on the sample surface, and cracks appeared in the region near the boundary as marked by arrow
After solution treated at 1160 °C, surface morphology of the alloy crept for different time. (a) After crept for 60 h, a few slipping traces appeared within the different grains, (b) after crept for 80 h, significant amount of the slipping traces appeared in the surface of the sample
After 1120 °C HIP alloy was solution treated at 1160 °C and twice aged, the morphology of the alloy crept for different time under the applied stress of 1034 MPa at 650 °C was shown in Fig. 37. The direction of the applied stress was marked by arrows, after the alloy was crept for 60 h, the morphology of the slipping traces on the sample surface was shown in Fig. 37(a), which displayed the feature of the single orientation slipping appearing within the different grains. And the intersected of the slipping traces appeared in the boundary region as marked by arrow in Fig. 37(a), which indicated that the boundary may hinder the slipping of the traces to change their direction. When crept for 80 h, the quantities of the slipping traces on the sample surface increased obviously, as shown in Fig. 37(b), and some white blocky carbide particles were precipitated within the grains.
After solution treated at 1160 °C and twice aged, the surface morphology of the alloy crept up to rupture under the applied stress of 1034 MPa at 650 °C was shown in Fig. 38. As the creep went on, the quantities of the slipping traces increases gradually (the direction of the applied stress shown in Fig. 38(a), which may bring out the stress concentration to promote the initiation of the micro-cracks along the boundary which was vertical to the stress axis as marked by the letter A and B in Fig. 38(a). In the other located region, the morphology of the crack initiation was marked by letter C in Fig. 38(b), the micro-cracks displayed the non-smooth surface as marked by arrow, and the white carbide particle was located in the crack, it indicated that the carbide particles precipitated along the boundary may restrain the cracks propagating along the boundaries to enhance the creep resistance of the alloy.
Cracks initiated and propagated along the boundary. (a) Crack initiated along the boundaries vertical to the stress axis, (b) crack propagated along the boundaries as marked by arrow
After the alloy crept up to fracture, the morphology of the sample polished and eroded was shown in Fig. 39. Some carbide particles were located in the boundaries as shown in Fig. 39(a), which may hinder the slipping of the dislocation for enhancing the creep resistance of the alloy. Moreover, the unsmooth surface of the cracks appeared in the fracture regions as marked by white arrow in Fig. 39(a). However, when no carbide particles were precipitated along the boundaries, the crack after the alloy crept up to fracture displayed the smooth surface as marked with the letter D and E in Fig. 39(b).
It may be thought by analysis that, although the carbide particles may hinder the dislocations movement for improving the creep resistance of alloy, the carbides located in the regions near the boundaries may bring about the stress concentration to promote the initiation and propagation of the cracks along the boundary as marked with the arrow in Fig. 39(a). Therefore, the fracture displayed the non-smooth surface due to the pinning effect of the carbide particles precipitated along the boundaries to restrain the boundaries slipping during creep. Though the carbide particles precipitated along the boundaries can improve the cohesive strength of the boundaries, the micro-cracks are still initiated and propagated along the boundaries, which suggests that the boundaries are still the weaker regions for causing fracture of the alloy during creep.
After solution treated at 1160 °C, surface morphology of the alloy crept up to fracture. (a) Carbide particles near the crack along the boundary marked by arrow, (b) morphology of cracks propagated along the boundary marked by arrow
After solution treated at 1165 °C, surface morphology of the alloy crept for 9 h up to fracture. (a) Crack initiated along the boundary as marked by arrow, (b) cracks propagated along the boundary as marked by arrow
After solution treated at 1165 °C and aged, the surface morphology of the alloy crept for 9 h up to rupture under the applied stress of 1034 MPa at 650 °C was shown in Fig. 40. A few slipping trace appeared only on the surface of the alloy, and some micro-cracks were initiated along the boundaries vertical to the applied stress axis, as marked by arrow in the Fig. 40(a). As the creep went on, the morphology of the micro-crack propagated along the boundary was shown in Fig. 40(b), in which the fracture of the alloy displayed the smooth surface. It may be deduced according to the feature of the smooth fracture that the carbide films precipitated along the boundaries has an important effect on decreasing the stress fracture properties of the alloy. The carbide films were formed along the boundaries during heat treated, which reduced the cohesive strength between the grains. Therefore, the micro-crack was firstly initiated along the boundaries with the carbide films, and propagated along the interface between the carbide films and grains, which resulted in the formation of the smooth surface on the fracture, and decreased to a great extent the creep properties of the alloy.
After the alloy was crept for 9 h up to rupture under the applied stress of 1034 MPa at 650 °C, the surface morphology after the sample was polished and eroded was shown in Fig. 41. The carbide films were continuously formed along the boundaries as marked with the long arrow in Fig. 41(a), the direction of the applied stress was marked by arrow, the micro-crack was initiated along the carbide film, as marked by shorter arrow in Fig. 41(a). As the creep went on, the morphology of the crack propagated along the boundaries was shown in Fig. 41(b), the fracture after the crack was propagated displayed the smooth surface, and the white carbide film was reserved between the tearing grains marked by arrow in Fig. 41(b), which displayed an obvious feature of the intergranular fracture of the alloy during creep. It can be thought by analysis that the carbide films precipitated along the boundaries, during heat treated, possessed the hard and brittle features and weakened the cohesive strength between the grains. Therefore, the micro-crack was firstly initiated along the carbide films and propagated along the interface between the grains and carbide films, which resulted in the formation of the smooth surface on the fracture, so the alloy had the lower toughness and shorter creep lifetime. Moreover, it was identified by means of composition analysis under SEM/EDS that the elements Nb, Ti, C and O were rich in the white particles on the surface of the samples, as shown in Fig. 37, Fig. 38 and Fig. 40, respectively, therefore, it is thought that the white particles on the surface of the samples are the oxides of the elements Nb, Ti and C.
After solution treated at 1165 °C, surface morphology of the alloy crept for 9 h up to fracture. (a) Crack initialed along the boundary marked by arrow, (b) morphology of cracks propagated along the boundary marked by arrow.
After the 1180°C HIP alloy was solution treated at 1150 °C and cooled in oil bath at 120 °C, the morphologies of the alloy crept for 260 h up to rupture under the applied stress of 984 MPa at 650 °C were shown in Fig. 42. If the PPB region between the powder particles was regard as the grain boundaries as shown in Fig. 42(a), the grain boundaries after the alloy was crept up to rupture were still wider, and the ones were twisted into the irregular piece-like shape as marked by arrow in Fig. 42(a).
Microstructure of alloy after crept up to fracture under the applied stress of 984 MPa at 650 °C. (a) Wider grain boundaries broken into the irregular shape as marked by arrow, (b) traces with double orientations slipping feature appeared within the grain as marked by arrows, (c) finer particles precipitated along the slipping traces
Some irregular finer grains were formed in the boundary regions, and displaying a bigger difference in the grain sizes. Some coarser precipitates were precipitated in the boundaries region in which the creep resistance is lower due to the spareness of the finer phase. The severed deformation of the alloy occurred firstly in the boundary regions during high stress creep, which resulted in the boundaries broken into the irregular piece-like shape. At the same time of the severed deformation, the traces with double orientations slipping feature appeared within the grains as marked by arrows in Fig. 42(b), and some particles were precipitated in the boundaries region as marked by short arrow in Fig. 42(b). Moreover, the finer white particles were precipitated in the regions of the double orientations slipping traces as marked by arrows in Fig. 42(c), and the white particles were distinguished as the carbides containing the elements Nb, Ti and C by means of SEM/EDS composition analysis.
Microstructure after the molten salt cooled alloy crept up to fracture under the applied stress of 1034 MPa at 650 °C. (a) Traces of the double orientations slipping appeared within the grains, (b) magnified morphology of the slipping traces
After solution treated at 1150 °C, and cooled in molten salt at 583 °C, the morphology of the alloy crept for 67 h up to rupture under the applied stress of 1034 MPa at 650 °C was shown in Fig. 43. This indicated that the traces with the double orientations slipping feature appeared within the grain, and the various orientations of the slipping traces appeared in the different grains, thereinto, the directions of the thicker and fine traces were marked by the arrows, respectively, in Fig. 43(a). Moreover, the traces with the cross-slipping feature were marked by shorter arrow in Fig. 43(a).
After solution treated at various temperatures, the alloy had different creep properties due to the difference of microstructure as shown in Table 6. When solution treated at 1150 °C, the alloy possessed a uniform grain size and wider PPB regions between the grains. Moreover, some coarser precipitates were distributed along the PPB regions in which no fine -phase was precipitated in the regions near the coarser -phase, as shown in Fig. 9(a), the regions possessed a lower creep strength due to the cause of the -free phase zone. After the alloy was solution treated at 1160 °C and twice aged, the coarser precipitates along the boundary regions disappeared, the boundaries appeared obviously in between the grains. And the cohesive strength between the grains was obviously improved due to the pinning effect of the fine carbide particles, as shown in Fig. 10(b), therefore, the alloy displayed a better creep resistance and longer the lifetime.
After the 1120 °C HIP alloy was solution treated at 1160 °C and twice aged, the alloy was crept for 104 h up to fracture under the applied stress of 1034 MPa at 650 °C, the fracture after the alloy was crept up to rupture displayed the initiating and propagating feature of the cuneiform crack as marked by letters A and B in Fig. 38. The schematic diagram of the crack initiated along the triangle boundary is shown in Fig. 44, where σn is the normal stress applied on the boundary, L is the boundary length, h is the displacement of the cuneiform crack opening, α is the crack length, θ is the inclined angle of the adjacent boundaries.
Schematic diagram of the crack initiated along the triangle boundary
Under the action of the applied stress, significant amount of the activated dislocations are piled up the regions near the boundary to bring the stress concentration, which results in the initiation of the crack in the region near the triangle boundary, and the crack is propagated along the boundary as the creep goes on. Thereinto, the critical length (αC) of the instable crack propagated along the boundary can be expressed as follows (Yoo M. H., 1983).
Where, G is shearing modulus, ν is Poisson ratio, f is the crack propagating work, h is the displacement of the cuneiform crack opening. This indicates that critical length (αc) of the instable crack propagated along the boundary increases with the displacement of the crack opening, and is inversely to the crack opening work. Thereinto, the displacement of the crack opening increases with the creep time, which can be express as follows:
Where hw=(τ/β) is the max displacement of the crack opening, τ is the resolving shear stress component applied along the boundary, t is the time of the crack propagation, δB is the boundary thickness, B is the sticking coefficient of the boundary slipping, β is the material constant.
The Eq. (8) indicates that the displacement of the crack opening (h) increases with the time and length of crack propagation. When two cuneiform-like cracks on the same boundary are joined each other due to their propagation, the intergranular rupture of the alloy occurs to form the smooth surface on the fracture. The schematic diagram of two cuneiform-like cracks initiated and propagated along the boundary for promoting the occurrence of the intergranular fracture is shown in Fig. 45. If the carbide particles are dispersedly precipitated along the boundaries, the ones may restrain the boundaries slipping for improving the creep resistance of the alloy to form the non-smooth surface on the fracture, as marked by arrow in Fig. 38(b).
After solution treated at 1165 °C and twice aged, the grain size of the alloy increased obviously, and the carbide films were formed along the boundaries as shown in Fig. 11, which weakened the cohesive strength between the grains. Therefore, the cracks were easily initiated and propagated along the boundaries adjoined to the carbide films, which may sharply reduce the lifetime and plasticity of the alloy during creep.
Schematic diagram of the cuneiform-like cracks initiated and propagated along the boundary. (a) Triangle boundary, (b) initiation of the cuneiform-like crack, (c) propagation of the crack along the boundary
Because the boundaries and the carbide particles can effectively hinder the dislocation movement, and especially, the carbide particles can improve the cohesive strength between the grains and restrain the boundaries slipping during creep, therefore, it may be concluded that the carbide particles precipitated along the boundaries have an important effect on improving the creep resistance of the alloy. Although the carbide particles precipitated along the boundaries can improve the strength of the boundaries, the micro-cracks are still initiated and propagated along the boundaries, which suggests that the boundaries are still the weaker regions for causing fracture of the alloy during creep. And once, the carbide is continuously precipitated to form the film along the boundary, which may weaken the cohesive strength between the grains to damage the creep lifetimes of the alloy. The analysis is in agreement with the experimental results stated above.
When the alloy was solution treated at 1150 °C and cooled in oil bath at 120 °C, the carbon atoms were supernaturally dissolved in the matrix of the alloy due to quenching at lower temperature. The concentration supersaturation in the alloy promoted the carbon atoms for precipitating in the form of the fine carbide particles during creep under the applied higher tensile stress at 650 °C, in especially, the slipping trace regions support a bigger extruding stress for inducing the carbon atoms to precipitate in the form of the fine carbide particles along the slipping traces as shown in Fig. 42(c). This is thought to be a main reason of the fine carbides precipitated along the slipping traces.
On the other hand, when the alloy was solution treated at 1150 °C and cooled in molten salt at 583 °C, although the slipping traces appeared still in the matrix of the alloy during creep, no fine carbide particles were precipitated along the slipping traces, as shown in Fig. 43, due to the concentration supersaturation of the carbon atoms in the matrix is lower than the one of the alloy cooled in oil bath at 120 °C.
By means of hot isostatic pressing and heat treated at different temperatures, creep curves measurement and microstructure observation, an investigation had been made into the influence of hot isostatic pressing and heat treatment on the microstructure and creep behaviors of FGH95 nickel-base superalloy. Moreover, the deformation and fracture mechanisms of the alloy were discussed. The conclusions were mainly listed as follows:
When the alloy was hot isostatic pressed below the dissolving temperature of γ′ phase, as the HIP temperature increased, the size and amount of primary coarse γ′ phase decreased gradually in the PPB regions, and the size of the grains was equal to the one in the previous powder particles. With the HIP temperature increased to 1180°C, the coarse γ′ phase in the PPB was completely dissolved, and the grain of the alloy grew up obviously.
When the solution temperature was lower than the dissolving temperature of phase, after solution treated at 1140 °C, finer phase was dispersedly precipitated within the grains, and some coarser precipitates were distributed in the wider boundary regions where appeared the depleted zone of the fine -phase. With the solution temperature increased, the amounts of the coarser phase and the zone of -free phase decreased gradually.
After solution temperature increased to 1160 °C, the coarser phase in the alloy was fully dissolved, the fine secondary phase with high volume fraction was dispersedly distributed within the grains, and the particles of (Nb, Ti)C carbide were precipitated along the boundaries. When the alloy was solution treated at 1165 °C, the size of the grains was obviously grown up, and the carbides were continuously precipitated to form the films along the boundaries.
During long term aging in the ranges of 450 °C and 550 °C, no obvious change in the grain size was detected in the alloy as the aging time prolonged, but the phase grew up slightly. With the aging time prolonging, the lattice parameters of the and phases increases slightly, but the misfit of phases decreased slightly.
Under the applied stress of 1034 MPa at 650 °C, the solution treated alloy cooled in molten salt displayed a better creep resistance. In the ranges of the applied temperatures and stresses, the creep activation energy of the alloy was measured to be Q = 590.320 kJ/mol.
The deformation mechanisms of the alloy during creep were the twinning, dislocations by-passing or shearing into the phase. The <110> super-dislocations shearing into the phase may be decomposed to form the configuration of (1/3)<112> super-Shockleys partial plus stacking fault.
During creep, the deformed features of the solution treated alloy cooled in oil bath was that the double orientation slipping of dislocations were activated, and the fine carbide particles were precipitated along the regions of the slipping traces. And the depleted zone of the fine phase was broken into the irregular piece-like shape due to the severe plastic deformation.
The deformed features of the alloy treated in molten salt were that the twinning and dislocation tangles were activated in the matrix of the alloy. Thereinto, the fact that the particles-like carbides were dispersedly precipitated within the grains and along the boundary might effectively restrain the dislocation slipping and hinder the dislocations movement, which is one important factor of the alloy possessing the better creep resistance and the longer creep lifetime.
In the later stage of creep, the slipping traces with the single or double orientations features appeared on the surface of the alloy. As the creep went on, the amount of the slipping traces increased to bring about the stress concentration, which might promote the initiation and propagation of the micro-cracks along the boundaries, this was thought to be the main fracture mechanism of the alloy during creep.
Most individuals seeking advice about weight management issues approach healthcare professionals with the intention of improving their health, well-being, and appearance by losing weight and altering their metabolic function. Typically, first responders, be they parents, nutritionists, physicians, nurses, coaches, or personal trainers, offer non-integrative behavioral approaches to weight management in the form of prescriptive diet plans. In the absence of underlying pathology or compulsions that may drive dysfunctional eating behaviors, such simple solutions may be adequate. But when the origin of an individual’s desire and efforts to lose weight resides in underlying eating pathology, purely behavioral solutions can mask potentially life-threatening dysfunctions. Sometimes hiding in plain sight, otherwise benign disordered eating (DE) habits may ultimately take on an element of compulsivity, leading to chronic illnesses such as heart disease, diabetes, and/or as heart disease or diabetes and/or the eventual onset of a life-threatening clinical eating disorder (ED) in genetically susceptible individuals. It is for the astute and intentional diagnostician to sniff out, intuit, or otherwise identify the potential for pathological origins within a constellation of seemingly benign weight management dysfunctions by determining whether a patient’s desire for weight change is based on healthful, autonomous discretion or on the dictates of a tyrannical and potentially lethal eating disorder. Overlooking the nature and severity of DE behaviors or warning signs of a clinical ED can carry dire consequences. By partnering with the patient to affirm or negate a diagnostic hunch, the clinician informs appropriate treatment, promoting disease prevention or the achievement of a timely and sustainable problem resolution.
The nature and quality of eating patterns reside along a continuum that includes eating behaviors and self-care. Along the span of this continuum, healthy eating patterns may evolve into disordered eating patterns, potentially leading to deadly ED in genetically susceptible individuals. At one end of the continuum, healthy eating behaviors mark the achievement of a fit and effectively functional body capable of sustaining its own ideal set point weight through healthy eating and self-care (Figure 1). At the opposite end of the continuum, life-threatening ED represent the tip of an underlying emotional, physiological, and neurobiological iceberg. A clinical ED marks the fragmentation or loss of the core self and, with it, the patient’s lost capacity for self-regulation, self-trust, self-esteem, and self-care. Feelings of guilt and of shame typically foster secrecy, denial, and reluctance to seek or sustain care. Ambiguous and inconstant DE behaviors exist somewhere between the polarities of healthy eating and potentially life-threatening ED. Some DE patterns that may appear to be pathological may actually represent variations of normal and benign behaviors, being shared by otherwise healthy eaters. As an example, healthy eaters who choose to indulge excessively in the lavish offerings of dessert buffets, typically consider such an opportunity benign and not-to-be-missed, guilt-free, gastronomic sensory and aesthetic delight.
Along the continuum of healthy eating, the evolution of healthy eating to DE, and from DE to a clinical ED, represents a journey of developing pathology; likewise, recovery from eating pathology represents a journey in the opposite direction, towards developing health.
The evolution of dysfunction along the eating continuum is influenced by the melding of eating lifestyle with genetic propensities, inborn temperament, value systems, developmental life stages, and self-concept issues culminating in body image perception and preoccupations. The confluence of normal child development and early bodily maturation, environmental triggers gleaned through social media, attitudes in the home, and peer pressure can lead to high-risk responses in teens and young adults which might include dieting, excessive exercise, abuse of diet pills, purging, food restriction, or eating only when hungry/skipping meals. In response to a societal disconnect between biology and culture, as girls grow older, they are expected to grow smaller. Studies show that 42% of first to third grade girls wish to be thinner [1] and, in the attempt to change one’s natural body shape and size, 81% of 10-year-old girls feel better about themselves when they are dieting [2], creating a greater risk of obesity in adulthood. 25% of American men and 45% of American women are dieting on any given day [3], giving rise to cravings, malnutrition, metabolic dysfunction, and increased overweight once the individual “falls off the dieting wagon.” It has been reported that 35% of “occasional dieters” progress to pathological dieting and DE, and as many as 25% progress to partial or full syndrome eating disorders [4].
It is for the perceptive and responsive diagnostician to determine at what point, along the healthy eating continuum, eating behaviors cross the thin line between normal eating and various forms and degrees of eating pathology and compulsions. A patient of mine with binge eating disorder (BED) assumed that she was in “remission” until, “It suddenly occurred to me that this was not just about having an extra piece of pie…this was starting to feel like an all-out binge.” Pathology exists when food, exercise, or the abuse of dietary substances is used to resolve emotional problems or in response to posttraumatic stress; when food serves purposes other than satiating hunger, fueling the body, or sociability; when the act of eating evokes guilt and fear; when eating habits become inflexible, immoderate, imbalanced, and excessive; when compulsive thoughts, such as calculating caloric intake, and preoccupations about becoming fat preclude other thinking, interfering with learning and normal daily function; and when attitudes and beliefs about food and eating are misguided.
When a child who is a healthy eater runs out the door without breakfast because (s)he is late for the school bus, she will surely want to snack during the morning or eat a bigger lunch to make up for the calories lost. The child with an ED who runs out the door without breakfast does so because (s)he would otherwise be wracked with fear, anxiety, and guilt that the calories she might ingest would make her fat. This child would feel compelled to calculate how many calories (s)he can allow herself to ingest throughout the rest of the day so as not to gain a single ounce.
In the face of weight management issues, information that patients communicate to clinicians is often not accurate or comprehensive; and what patients request of health professionals is frequently not what they need. Falling between the cracks of accurate assessment and appropriate treatment, DE and potentially lethal ED all too frequently remain misunderstood, underdiagnosed, and undertreated [5] by myopic, underinformed physicians, psychotherapists, and nutritionists who lack an integrative “big picture” perspective of the possible existence of underlying, co-occurring diagnoses that need to be revealed, monitored, and treated. The diagnosis of a clinical ED remains elusive, with more than one half of all cases remaining undetected [6].
A patient of mine had been treated by a psychiatrist for 7 years, during which time, she never told him that she vomited 30 times a day, for fear of his finding her “disgusting” and therefore refusing to treat her. This physician failed to “read between the lines” of her symptoms and to intuit and investigate the presence of an ED from the constellation of personality traits, temperament, and behaviors that she did reveal to him. Through our work together, she came to understand her problem and her self. The quality of our relationship led to the improvement of all her significant relationships—with food, with her self, and with others.
The unique requirements of weight management issues demand the uniquely specialized skills of an informed, intentional, and intuitive diagnostician, capable of hearing what has not yet been spoken. Within the context of initial history-taking, the nature of weight-related dysfunctions is likely to remain elusive in the absence of an active probe for problem origins in pathology. The diagnostician’s enlightened line of questioning will substantiate, or negate, such a presence, revealing the subtleties of distinctions between the nature of eating behaviors as they reside along the healthy eating continuum. When eating behaviors do appear to cross the line into pathology, the diagnostician takes on the role of crisis interventionist, through immediate responsiveness to the needs of the moment through an investigation into the past. Because the patient’s quality of life may depend upon assessment accuracy, the obligation for first responders or clinicians to detect, explore, and interpret issues yet to be unearthed becomes a unique challenge within the confines of the limited timeframe of a single-session weight management consultation. On high alert for potentially unseen issues, the proactive first responder must be prepared to offer psychoeducation, a plan of action, and, where appropriate, referrals to collaborating experts and/or higher-level treatment milieus. If it looks like a duck and acts like a duck, it is for the responsible practitioner to treat it like a duck, unless proven otherwise, even before the assignment of a definitive diagnosis.
Esther was a 29-year-old woman who came to treatment for depression, poor self-esteem, and relationship problems. In response to her description of her college days during which she spoke of herself as perfectionistic, highly compulsive, anxious, and depressed, I chose to wonder aloud if she had ever struggled with an ED or other eating-related issues. “My God!” she responded. “How did you know? I have never told a soul!” By understanding the nature of her personality structure and recognizing characteristics of her emotional functioning, I was able to intuit and surmise the possible existence of a past ED, which upon inquiry, I discovered had yet to be fully resolved. Following the tenet of John Muir, “If we try to pick out anything by itself, we find it hitched to everything else in the universe.” By following my hunch and having made this discovery, I was able to launch the ED treatment process immediately.
Where the clinician’s depth and breadth of understanding of weight management issues is limited and non-integrative, the patient is liable to leave the treatment office with a diet plan in hand, yet without a practicable and sustainable solution for pressing weight-related concerns. More significantly, the loss of a timely and poignant opportunity for the patient to discover unknown problem origins and aspects of self and personality represents a lost opportunity for self-reflection, self-awareness, and self-integration, all leading to healthier lifestyle choices.
Factors further clouding the differential diagnosis of eating-related pathology include the elusive assessment of evolving and ambiguous DE patterns, and the widespread misunderstanding and erroneous information surrounding the issues of eating, weight control, and ED. In asking my ED patients what triggered the onset of their ED, a frequently heard response is that they “started to diet in an effort to lose weight, and ultimately found themselves feeling compelled to eat ‘healthier and healthier.’” The term “healthy eating” not infrequently becomes a euphemism for food restriction, or ‘clean eating,’ potentially triggering the onset of orthorexia in ED individuals. Further complicating the diagnostic process, ED symptoms vary appreciably from patient to patient, with; every ED a ‘thumbprint.’ Fully half of the ED population suffers from the difficult-to-identify condition called “other specified feeding or eating disorder” (OSFED), previously known as “eating disorder not otherwise specified” (EDNOS) [7]. OSFED describes atypical AN (without low weight) and atypical BN or BED (with lower frequency of behaviors, purging disorder, and/or night eating syndrome). Patients who do not meet the strict diagnostic cutoffs for full criteria for AN and BN often remain undiagnosed despite the seriousness of their illness, foregoing or delaying necessary treatment [8]. Patients with clinical ED are often reticent, within a diagnostic interview, to divulge a stigmatized disorder. A recent discovery, important for its potential to reduce the degree of stigma that is associated with ED nondisclosure revealed that 20% of the neurobiology of AN could be derived from metabolic genes (possibly activated by a state of starvation) [9].
Consider the following examples of decoys to the recogntion and understanding of ED:
People mistakenly assume that AN is easy to spot, believing that “all anorexic individuals are exceedingly thin.”
It is a commonly believed misconception that an AN individual’s full restoration of weight marks a full recovery from an ED. In fact, ED recovery is marked by neurobiological, emotional, cognitive, and behavioral changes that lead to the reintegration of the individual’s core self and the normalization of body fat mass and sexual hormones, which have a widespread impact on the body and multiple pathways in the brain [10].
Despite the widely accepted misconception that ED represent an incurable “life sentence,” about half of those individuals with AN or BN attain a full recovery, 30% achieve a partial recovery, and 20% show no substantial improvement [11].
Though early detection of ED warning signs is pivotal in disease prevention and/or promoting a timely recovery, the medical community has been known to overlook opportunities for early disease recognition.
Because signs of an ED rarely appear in blood tests until advanced stages of disease, normal test results in early stages of ED are often misconstrued to represent a “clean bill of health.”
Pediatricians frequently overlook the significance of precipitous weight loss in children when their numbers fall within the range of normal on the growth charts, precluding early disease prevention.
Gynecologists regularly prescribe birth control pills for ED patients with amenorrhea, erroneously assuming that hormone replacement will counteract or reverse bone loss and improve reproductive functionality [12].
Cardiologists, in the face of a co-occurring activity disorder that takes the form of excessive and compulsive exercise, frequently miss an ED diagnosis by attributing a low heart rate and amenorrhea to “healthy athleticism.”
When driven by the fear of weight gain, individuals who suffer from ED, co-occurring anxiety, and obsessive-compulsive disorders (OCD), are particularly susceptible to developing activity disorders. Also known as anorexia athletica, exercise bulimia, or exercise addiction, ED individuals engage in such compensatory compulsions with the intention of burning calories. Between 40% and 80% of AN patients are prone to excessive exercise in their efforts to avoid putting on weight [13].
Binny ran 10 miles a day, followed by 2 hours of working out. She ate no more than 750 calories a day while training regularly for countrywide marathon races. One 26.2-mile event landed her in a hospital, where her legs swelled and she required an emergency blood transfusion. Because her eating and running regime provided her with a sense of being “alive,” upon release from the hospital she felt incapable of curtailing the compulsive behaviors that threatened her life. Her emergency room doctors attributed her blood disorder and amenorrhea to her athleticism, failing to recognize both conditions as signs of ED pathology.
“The beginning of wisdom is to call things by their right names.”
Chinese Proverb
Barring origins in genetic, metabolic, or hormonal dysfunctions, weight management dysfunctions typically originate in an individual’s unhealthy relationship with food, leading to a disordered eating lifestyle. The term “weight management” describes the techniques and physiological processes that contribute to attaining and maintaining an individual’s ideal weight. Healthy weight management techniques encompass long-term lifestyle strategies promoting healthy eating and daily physical activity, fostering sustainable change and well-being. In contrast, unhealthy weight management strategies, lacking an integrative treatment perspective, fail to achieve sustainable weight goals. Examples of unhealthy weight management include dieting, skipping meals, food restriction, eating only when hungry, and forms of purging that may include vomiting, spitting, compulsive exercise, and the abuse of laxatives, diet pills, and diuretics. Multiple studies have found that dieting for purposes of weight management is associated with greater weight gain and increased rates of binge eating in both boys and girls [14].
Healthy eating is guilt-free, balanced, and fearless eating, with flexibility in accommodating the parameters of the dining moment. Healthy eating includes three meals daily, each including all the nutritionally-dense food groups, as well as snacks. There are no bad foods. What is bad is extremism, compulsivity, and unhealthy attitudes about food and eating, i.e., how we feel about what we eat. What is worse than eating Oreos is never eating Oreos, as forbidding a child to eat Oreos can ultimately lead to sneaking, hiding, or stealing food. Healthy eating results in healthy weight maintenance naturally, through the inherent wisdom of a body which, when fed healthfully, and having reached its ideal set point weight, will predictably sustain that appropriate weight-to-height ratio indefinitely, through its own natural functions. In most cases, the healthfully fed and exercised body with a healthfully functioning metabolism needs no help from arbitrarily devised dietary interventions to sustain a healthy weight throughout one’s lifetime.
At times we are all a little eating dysfunctional. Bombarded by nutritional research and food fads in an age of pervasive social media, “normal” eating (i.e., eating all food groups, including moderate amounts of processed foods and sugar) is no longer considered to be “healthy” eating. Certain patterns of DE, in light of their prevalence, are becoming increasingly difficult to recognize as pathological. A study found that nearly 91% of female college students use dieting as a weight control mechanism [15], putting many of them at risk to develop a clinical ED. The prevalence of AN and BN is relatively rare among the general population [16], affecting 1% to 4% of adolescents and young adult women [17]; DE, however, which has been defined as a psychological illness, [18], and the misguided attitudes about food and weight management that lead to them, are widespread and prevalent among all age groups. The consequences of DE, which include rampant dieting and body image obsessions, can be devastating; though not the cause of eating disorders, they are often precursors to their onset. Occasional DE, as well as certain behaviors that manifest themselves in clinical ED, are not in themselves abnormal. Differentiating normalcy from pathology is best accomplished by assessing the patient’s attitudes towards food and eating. Benign forms of DE are fully discretionary, dependent upon the individual’s capacity for self-awareness, self-regulation, and choice-making through their own free will.
Picky eating syndrome is a DE pattern that originates in early childhood. Typically indicative of a sensory processing disorder (SPD), food choices tend to be limited to bland “white” foods (containing sugar and flour) of specific textures. Picky eating often results in overweight. Studies show that 40% of overweight girls and 37% of overweight boys are teased/bullied about their weight by peers or family members and that traumatic experiences such as these during the formative years are predictive of weight gain, binge eating, and extreme weight control measures [14]. DE patients have experienced a higher frequency of traumas (childhood adversities in particular), especially in circumstances related to childhood obesity [18]. Childhood picky eaters who do not grow out of the condition and who become malnourished because of the limited variety of foods that they eat qualify for a diagnosis of avoidant restrictive food intake disorder (ARFID), a condition which typically extends into adulthood.
Rising obesity rates, the result of DE lifestyles, clinical ED, and hereditary factors have become a major concern worldwide, especially in North America, where more than 2 in 3 adults (70.2%) are considered to be overweight or obese [19]. Obesity, DE, and ED are considered major health problems among adolescents because of their increasing prevalence in this age group and their potentially serious physical and psychosocial consequences [20, 21]. Dieting and unhealthful weight control practices lead to eating psychopathology in DE, ED, and obesity, particularly when associated with variables such as personality characteristics like perfectionism or anxious temperament, a microbial imbalance, or a genetic predisposition to be influenced by an obesogenic environment. Other possible causal mechanisms linking dieting to later problems are neurophysiological mechanisms associated with food restriction (e.g., effects on neurotransmitters that could increase risks for either obesity or ED via influences on food regulation processes) [21]. In the face of a genetic, hormonal, or metabolic predisposition to obesity, following a healthy lifestyle in which the individual learns to eat differently, not less, may counteract gene-related risks. Individuals with obesity and co-occurring eating disorders are at higher risk for several medical and psychosocial complications than individuals with either condition alone [22]. Because obesity may become a precursor to an ED, and vice versa, collaborative exchanges of experiences and specialized knowledge between healthcare professionals working in the fields of obesity and eating disorders are essential [22].
ED onset is an integrative process, stimulated by contributing factors that include genetics, neurobiological vulnerability, physiology, co-occurring mood and emotional disorders, past trauma, and familial and cultural influence. For the individual with AN, BN, BED, or OSFED, behavioral and emotionally based compulsions become etched in neuronal pathways, impacting the structure and function of the brain. ED patients consider their disorder to be a survival tool, an adaptive coping mechanism that they cannot live without. Underlying the clinical ED is the primordial fear of becoming fat. For the ED individual, food is “fattening,” and fat, the enemy of a body that can never be too thin, becomes a “feeling,” a sensation, a perception. ED are self-regulatory dysfunctions. All types of ED share a great number of symptoms and issues, as aspects of the same disease syndrome; however, ED victims vary in the basic structure and function of personality and emotional makeup, and treatment needs to vary, reflecting these differences. BN behaviors tend to be marked by impulsivity, followed by “undoing” practices, such as various forms of purging. AN behaviors reflect behaviors and attitudes marked by rigidity, restriction, and containment.
Though ED are essentially not about food, their successful healing depends upon the resumption and maintenance of a healthy weight and relationship with food [22]. ED are disorders of the core self, characterized by diminished self-control, self-regulation, self-attunement, self-trust, self-agency, self-reliance, self-perception, self-sensing, and self-worth. A full ED recovery is defined by several factors, including the return of a healthy and sustainable eating lifestyle and relationship with food; the reintegration of the mind, brain, and body, reconstituting a fully integrated self; the arrival at one’s set point weight; and a full restoration of healthful physiological functioning leading to the natural resumption of menstruation. Observable weight change is typically the factor that brings ED patients to treatment. Once in treatment, additional predictive assessment factors include the ED mentality, marked by perfectionistic, obsessive, black-and-white thinking, cognitive rigidity, the compulsive quality of behaviors, the intolerance of uncertainty, and the fear of gaining weight and becoming fat. It is important to note that many individuals with various forms of ED maintain a normal weight, a phenomenon observed particularly in BN and BED because of “weight cycling,” where patients alternate between starvation and gorging or compulsive bingeing followed by compensatory forms of purging.
ED do not stand still; any ED that is not in the process of healing is becoming increasingly entrenched in the brain and nervous system, destroying life quality, if not taking lives.
Pamela suffers from BED. She is not overweight, as might be expected, the result of yo-yoing between gorging and starvation for days on end. Aware of her illness and her need for professional help, she spoke of her frustration in not having been able to get the attention of her family or the healthcare professional community, who refused to acknowledge that she had an ED because her weight was “normal.” Having been denied treatment, compassion, much needed support and attention, upon arrival at my office, she spoke of feeling depressed, isolated, and hopeless.
In looking beyond simple, nonlinear solutions for weight management problems, the single-session diagnostic assessment needs to become a discovery process, uncovering the possibility of underlying sources of symptoms in eating pathology or past trauma. In establishing a direction for future treatment, the assessment of DE or ED requires treatment tools of subtle refinement that investigate the patient’s current disease and recovery status, as well as the history of past efforts to heal. Following Abraham Maslow’s hierarchy of human needs [23], the weight management diagnostician attends to the needs of the patient in a sequence designed first to save lives through medical stabilization, then to remediate life quality through emotional stabilization. For ED patients, the need to refeed body and brain becomes a first priority in minimizing the patient’s physiological risk, augmenting receptivity to the therapeutic process and to an environment conducive to the use of medication. Physiological or emotional instability that could result in self-harm or death demands an immediate referral to higher levels of care.
Psychoeducation lies at the core of the patient’s engagement in the healing process as a mainstay of support, rectifying distorted attitudes and cognitive belief systems; introducing a deeper understanding of weight management problems and their implications; opening the patient’s eyes and mind to insights and self-reflection; and creating the patient’s sense of trust in the clinician and the clinical process as well as hope for a full recovery. Psychoeducation emphasizes the importance of maintaining a healthy eating and activity lifestyle to insure short-term and long-term goals for unified mind, brain, and body health. A psychoeducational explanation of the set point weight theory becomes critical in assuaging the AN patient’s fear of excessive weight gain during refeeding by describing the wisdom of the human body, which, once having restored optimized bodily functioning, will sustain a constant weight. The AN patient also learns to anticipate that the natural course of ED recovery will involve countless trials and setbacks in regaining lost weight, with the recognition that every movement representing progress or regression represents “grist for the learning mill.” Patients and clinicians need to anticipate the emergence of previously buried feelings during the refeeding process, evoking states of psychophysiological fear and/or emotional distress. Psychoeducation is also of great benefit for the parents of eating dysfunctional children or young adults living at home. Through family therapy, parents become knowledgeable about the complexities and risks involved with their child’s eating dysfunctions and the urgent need for total healing. As advocates for their ED child’s recovery, skillfully coached parents become, and remain, effective agents for positive change throughout an ever-changing treatment and recovery landscape.
Aside from psychoeducation, the work of the initial diagnostic session or series of early sessions needs to be richly flecked with trust-building and relationship building, along with an action plan, devised even prior to the development of a definitive diagnosis. The assignment of tasks such as journaling, requests for on-going open and honest personal feedback, and contingency contracts all foster the patient’s learning, trust, and treatment engagement. The initial session may include professional referrals to prospective members of a treatment team, as needed, initiating an integrative treatment process. Evoking motivation for change, the quality of the initial patient/clinician connection awakens the patient’s recognition and acceptance of oneself, of co-occurring diagnoses that may require attention, and of the need for commitment to a treatment process that can accomplish full and sustainable healing. Aside from identifying current behaviors, levels of function, and past treatment and recovery efforts, history-taking needs to assess the patient’s internal strengths, external resources, environmental influences and mood dysfunctions, always with attention to the possibility of past or current trauma. Studies that include partial or subclinical forms of posttraumatic stress disorder (PTSD) show that well over half of individuals with BN have PTSD or significant PTSD symptoms [24]. In light of the prevalence of trauma in the background of ED patients, trauma investigation needs to become a central focus within the initial inquiry. The diagnostician does well to become self-aware of personal propensities towards countertransference responses or cultural biases (reflecting weightism). The latter are forces which might preclude the clinician’s recognition that many overweight and genetically large individuals who eat healthfully and exercise regularly are physically fit, healthy, and strong.
A heavy burden of medical comorbidities across multiple body systems, attributable to both the malnutrition of AN and the purging behaviors of BN, contribute to the high mortality rates of these illnesses [25]. Restoration of weight and nutritional status are key elements in the treatment of AN [26]. Nutritional and medical treatment of extreme undernutrition present two very complex and conflicting tasks: the need to avoid “refeeding syndrome” caused by a too fast correction of malnutrition and “underfeeding” caused by a too cautious refeeding [27]. Metabolic, endocrinological, and gastronomic consequences that may develop during the refeeding process for ED individuals in recovery need to be understood and addressed in treatment. For the most part, but not in all cases, adverse consequences are reversible with recovery:
Efforts of AN patients to restore weight within the refeeding process risk derailment due to the irrational fear that normal eating behaviors will ultimately lead to overweight or obesity following weight restoration. A psychoeducational discussion of optimal set point weight can potentially diminish fears through the knowledge that the body will maintain this optimized weight by using the same energy intake as had been needed for weight restoration. With metabolic normalization and biological functions coming back on line after their dormancy during weight restriction, the extra energy used previously for weight gain becomes expended on usual day-to-day functions [28].
During initial phases of weight restoration, particularly for restricting ED patients who begin treatment at lower weights, the metabolic rate may overshoot normal levels in a “hypermetabolic” phase in which patients easily lose weight and need to eat an even larger amount of food to gain and sustain weight. This phenomenon is due to increased diet-induced thermogenesis (with calories dissipated as heat) as well as a variety of neuroendocrine alterations [26]. Despite the urgent need for the body to use restored energy efficaciously to replenish fat reserves and repair tissues in the early weeks and months of refeeding following prolonged semi-starvation, metabolic function may not normalize for 3–6 months following weight restoration [28].
Target weights offer invalid markers of metabolic normalization, providing false indicators of ED recovery progress. It has been shown that ED patients forced to gain weight in hospital settings typically plan to lose it upon discharge [29]. Attempting to attain a target weight that is not reflective of the body’s self-determined set point is predictive of a poor long-term prognosis. Establishing and maintaining a goal-weight range is but one of many integrative physiological, neurobiological, nutritional, cognitive, emotional, and environmental factors contributing to a complete and sustainable ED recovery.
Weight gain from edema (swelling or bloating) can be caused by hormonal changes brought on by starvation. AN can lead to kidney-related issues that include severe electrolyte disturbances (hypokalemia, hypophosphatemia, etc.), nephrolithiasis, and alterations in water metabolism (with hyponatremia and edema.) Patients with the binge eating/purging subtype of AN are more likely to have kidney disorders, particularly electrolyte disturbances and volume depletion, than those with the restrictive subtype. ‘Refeeding edema’ occurs during the process of weight restoration [30].
Weight gain is a common side effect for people who take insulin to manage type 1 diabetes. The deliberate insulin underuse with type 1 diabetes for the purpose of controlling weight is known as diabulimia, a condition leading to a threefold increase in mortality risk [27].
Sarcopenia, the loss of muscle mass in moderate to severe AN is frequently overlooked by clinicians. Appropriate exercise is required to restore muscle mass and strength [31].
The strongest predictors of osteoporosis include low body weight and amenorrhea. Hormonally based bone softening and/or bone loss may never become restored to a state of normalcy, even following weight restoration.
AN has consistently been associated with reduced gray matter and white matter brain volumes. It remains unclear if gray matter alterations are present following recovery from AN [32].
Issues complicating the achievement of full weight restoration matter a great deal, as halting the process of weight restoration at much lower levels than needed for full and sustainable recovery increases the rate of treatment dropout [33] and relapse [34]. The restoration of nutrient status and weight needs to start slowly and gradually, accelerating as tolerated. The refeeding process needs to focus on modifying the disordered dietary patterns that AN patients commonly practice, which might include slow and irregular eating, vegetarianism, and the consumption of a restricted range of foods. Severely malnourished AN patients often need to be admitted to a hospital in order to receive more aggressive treatment, with extra care and monitoring required to prevent the occurrence of refeeding syndrome [26], a clinical complication involving kidney dysfunction [30]. A person may finally be considered “in remission” after maintaining a stable weight for a number of years and experiencing the natural resumption of mensuration, as well as other normal hormonal, metabolic, and gastroenterological processes. Sustained remission is marked by the return of the reintegrated core self, following its fragmentation by the encroachment of the ED “pseudo-self.”
Trauma in the form of sexual abuse occurs in 30–65% of women with ED [35]. The vast majority of women and men with AN, BN, and BED report a history of interpersonal trauma, with approximately one-third of women with BN meeting criteria for lifetime PTSD [36]. Unresolved trauma and/or PTSD can be an important perpetuating factor in the maintenance of ED symptoms [24]. DE and ED behaviors typically serve the patient as coping tools and distractions that numb the ongoing effects of traumatic memories. A trauma-focused approach to ED treatment facilitates the resolution of traumatic experience that lies at the root of behavioral, emotional, and neurobiological dysfunctions [37].
Unprocessed traumatic memories stored in the mid-brain region become recycled when triggered, creating undischarged energy in the nervous system. Any traumatic assault on, or insult to, the brain impairs brain integration. High stress levels that lead to an overactive amygdala and hippocampus suppress the activities of the prefrontal cortex, the thinking brain that helps to regulate emotions. Body dysmorphic disorder, a common occurrence in ED individuals traumatized by sexual abuse, represents a mind, brain, and body disconnection within the disparate nervous system. BN and AN pathology reflects the disintegration of the structure of the self within the distributed nervous system, resulting in the patient’s inaccurate sensing of self-based experience and perception of self. Psychosomatic expressions of traumatic experience are held as bodily sensations. ED heal in the same way as trauma heals, through the neurophysiological and neurobiological reintegration of the distributed nervous system, marking the return of the patient’s reintegrated core self [38]. Because traumatic memories are encoded subcortically, the process of healing ED that originate in trauma requires accessing, and gaining leverage within, the structural coding of the brain and nervous system.
Neurophysiological effects of past trauma that are revealed in the present become accessible and available for remediation [38]. Trauma resolution lies in creating a psychophysiological state associated with decreased adrenergic activity, decreased muscular neuromuscular arousal, and cognitive quieting [39]. The introduction of neurophysiological (sensorimotor) and neurobiological (interpersonal, attachment-based) treatment interventions into mainstream clinical treatment for ED increases exposure to mindful embodied movement experience, fostering mind, brain, and body connectivity. By stimulating integrative neuronal firing and synaptic activity, these “top-down” and “bottom-up” transactions enhance acuity in self-sensing, self-perception, and body image coherence, supporting the unification of the disparate self [38].
Trauma resolution becomes enhanced through a mindful, quality connection between the therapist and patient. Rapid resolution therapy (RRT) is a body-based talk therapy technique shown to alleviate negative effects of trauma and PTSD without requiring the patient to recollect painful memories [40]. Trauma resides in the limbic system and in the perceptual world within a neural network, which has sufficient functional boundary thresholds to largely “disintegrate” it from the rest of the nervous system. When negative feelings become dissociated or “split off” (as they do within the bulimic “pseudo-self”), the potential exists to reintegrate them through the patient’s connection with his/her best and resourceful self, through solution discovery, or rediscovery, both past and present [40]. Trauma resolution accesses neuroplasticity, through which neural networks that become lit-up at the same time as the neural network associated with the problem result in the problem’s loss of definition. This dynamic allows for a free flow of communication with the rest of the nervous system, as the brain reinterprets new combinations of neural connections to create meaning [40]. By creating connections within the distributed nervous system in the context of a trusted human relationship, the technique connects problems to solutions through consolidating memories of human strengths and resourcefulness.
For decades under the care of dietitians treating Charles’ disordered eating, dieting efforts failed, leaving him intractably obese at 100 pounds overweight. Still “feeling fat” following bariatric surgery, he came to therapy struggling to lose weight through old patterns of food restriction, leading to junk food binges. At the start of our treatment, history-taking revealed past trauma at the root of his current eating dysfunction. Since childhood, Charles’ family had been forced to flee a war-torn country with only a day’s notice, leaving their previously comfortable lives and extended family behind. Moving from country to country, they struggled to survive as refugees, grieving their losses and experiencing years of fear, hunger, and social isolation. In response to the combined work of ED recovery and trauma resolution, within the context of our quality therapy relationship, Charles began to experience a sense of grounded integration and selfhood, evidenced in his growing capacity for self-regulation and self-care. Within several months of treatment, he became a normalized healthy eater and started to exercise regularly. Change in his distorted body image perception as a fat person became apparent when he reported, “During an exercise class, I noticed myself in the mirror and thought, ‘Who is that normal weight person? Could that be me?’ Despite his weight management problems and eating compulsions, the healthy personality at his core had resurfaced through psychotherapy. In completing his treatment, he described himself as having become “a happy and gratified human being, enjoying a fulfilled existence.”
In some instances, a skilled practitioner may be capable of discovering trauma and shepherding the start of the trauma resolution process within the context of the initial time-extended single-session weight management assessment.
It was the sudden recurrence of compulsive bingeing habits that she’d assumed were “in remission” that brought this middle-aged woman with BED to treatment. “This is the one area of my life that has always remained just beyond my grasp…. all I know is that my hunger is insatiable.” In probing for the possibility of past trauma, our discussion during that initial session uncovered the source of her erratic eating compulsions in feelings of shame, emptiness, and emotional lability originating in early childhood trauma, a connection that had remained outside of her consciousness for 45 years. Her memories of neglect, alienation, and hunger that she suffered at the hands of her parents throughout her growing-up years, having been buried for decades within her limbic brain, now began to surface. Her brain’s rapidly firing neuronal connections brought forth immediate new insights into herself, her current feelings, and past emotions, evoking an enlightened clarity about her previously “incomprehensible, over-reactive responses” (sobbing tears, sleeplessness, irrational fears) in the face of certain types of stressful experiences throughout her life. She wept with relief and gratitude upon leaving, recognizing that significant and sustainable change for her would now be in the offing.
Weight management anomalies signify the patient’s unhealthy relationship with food, potentially giving rise to chronic disease or the onset of genetically predisposed clinical eating disorders and the metabolic, endocrinological, or gastrointestinal consequences that characterize them. Uncovering the origins of weight management dysfunctions in DE or ED pathology, in providing a direction for treatment, enhances the efficacy and sustainability of healing. From the perspective of a hammer, all things look like a nail; in assessing weight management issues through a purely behavioral lens, first responders who fail to investigate and probe an underlying emotional landscape are likely to recommend a non-integrative solution, such as dieting, missing the opportunity to address the full complement of impinging psychological, neurobiological, and neurophysiological factors that contribute to weight management dysfunctions. Dieting, and particularly the use of unhealthful weight control behaviors, increase risk for weight gain and later eating- and weight-related problems [21].
It is through a unique use of self within the diagnostic moment that the knowledgeable and informed first responder approaches the uniquely challenging arena of weight management. By “listening with a third ear,” the clinician with clear intention and exquisite sensitivity to implied and unspoken issues intuits and then skillfully addresses the possibility of a yet unknown and unnamed condition as part of a wider constellation of symptoms. The proactive diagnostician fosters the patient’s self-reflection, self-acceptance, and incentive to heal dysfunctions sustainably at their source; increased self-esteem and positive body image have been shown to be best achieved through self-acceptance rather than weight reduction [41].
Research justifies the need for long-term implementation of interventions that aim to simultaneously prevent the onset of obesity and ED through the prevention of dieting behaviors and the promotion of healthful eating and physical activity as ongoing lifestyle behaviors [21]. Research reveals that fitness center employees, ideally placed to observe clients who exhibit an addiction-like relationship with exercise in an effort to lose weight as part of an ED, require detailed guidelines for intervention, including ways to start conversations to this end [42]. Bottom line, first responders need to determine whether the patient’s desire for weight change is based on healthful choices and discretion, or on the dictates of pathological compulsions that underlie and drive dysfunctional eating behaviors. In either case, the diagnostician sets the stage for the patient’s immediate and compelling engagement in integrative treatment, creating the potential to save lives and promote life quality.
Heartfelt thanks, love, and gratitude to my husband Lou, who edited this chapter with devotion and enthusiasm.
The authors declare no conflict of interest.
DE | disordered eating |
ED | eating disorders, eating disordered |
BED | binge eating disorder |
OSFED | other specified feeding or eating disorder |
EDNOS | eating disorder not otherwise specified |
OCD | obsessive-compulsive disorder |
ARFID | avoidant restrictive food intake disorder |
SPD | sensory processing disorder |
PTSD | posttraumatic stress disorder |
RRT | rapid resolution therapy |
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