UNCLASSIFIED, UNC LASSIFIED| THE UNIVERSITY OF MICHIGAN, Ann Arbor, Mich. THE UNIVERSITY OF MICHIGAN, Ann Arbor, Mich. EFFECT OF CREEP-EXPOSURE ON MECHANICAL PROP- EFFECT OF CREEP-EXPOSURE ON MECHANICAL PROPERTIES OF RENE' 41: Part II, Structural ERTIES OF RENE' 41: Part II, Structural Studies, Surface Effects, and Re-Heat Treat- Studies, Surface Effects, and Re-Heat Treatment, by Jeremy V. Gluck and James W. ment, by Jeremy V. Gluck and James W. Freeman, November 1961. 103p. incl. figs. Freeman, November 1961. 103p. incl. figs. tables and refs. (Project 7381; Task 73810) tables and refs. (Project 7381; Task 73810) (ASD TR 61-73, Pt. II) (Contract AF 33(616)- (ASD TR 61-73, Pt. II) (Contract AF 33(616)6462) Unclassified report 64 Unclassified report The effect of creep-exposure on room temp- The effect of creep-exposure on room temperature mechanical properties of Rene' 41 erature mechanical properties of Rene' 41 was studied for temperatures of 1200-18000F was studied for temperatures of 1200-1800~F and times up'to 200Op.urs. Unstressed ex- and times up to 200 hours. Unstressed exposures were for as long. as 2012 hours at posures were for as long as 2012 hours at 17000F.'Thermally-induced structural 17000F. Thermally-induced structural changes reduced strength ana ductility after changes reduced strength and ductility after exposures at 1400-18000F. Reduced yield exposures at.]+00-18000F. Reduced yield UNCLASSIFIED| UNCLASSIFIED| (over ) over)'S'-r - - _- -- -- - _.- I - _ -- _ a strengt waUNCLASSIFIED UNCLASSIFIED strength was due to decreased volume fraction strength was due to decreased volume fraction of gamma prime and secondarily to an in- of gamma prime and secondarily to an increase in the particle size, Ductility was crease in the particle size. Ductility was reduced by formation of massive grain bound- reduced by formation of massive grain boundary carbides. Up to 15000F, creep caused ary carbides. Up to 15000F, creep caused strain hardening. and pauischinger effects. strain hardening and Bauschinger effects. Except for surface effects, damage was re- Except for surface effects, damage was restorable by re-heat treatment. Yield storable by re-heat treatment. Yield strength was restored by re-solution and re- strength was restored by re-solution and reaging to produce fine gamma prime. Complete I aging to produce fine gamma prime. Complete re-solution of carbides was required to re- re-solution of carbides was required to restore ultimate strength and ductility. store ultimate strength and ductility. Microcracking was not observed. Creep in- Microcracking was not observed. Creep induced intergranular surface cracking at duced intergranular surface cracking at 1200-13000F which reduced ductility. Sur- 1200-13000F which reduced ductility. Surface effects for exposures above 14000F were face effects for exposures above 14O0F were thermally induced. General principles were thermally induced. General principles were formulated for damage to properties of formulated for damage to properties of nickel-base alloys. UCLASSIF nickel-base alloys. UNCLAS 1 -C- - _ASS IFE _ UNC -AS S F IED - A y _ ~ __ __- __ __ __- __. __ _ _.... _. __...~. - -_ _~.. J- - - __- __ _ _ _ _._ __ _ __ _ _ _ _ _

UNCLASSIFIED UNCLASSIFIED THE UNIVERSITY OF MICHIGAN, Ann Arbor, Mich. THE UNIVERSITY OF MICHIGAN, Ann Arbor, M\ich. EFFECT OF CREEP-EXPOSURE ON MECHANICAL PROP- EFFECT OF CREEP-EXPOSURE ON MECHANICAL PROPERTIES OF RENE' 41: Part II, Structural ERTIES OF RENE' 41: Part II, Structural Studies, Surface Effects, and Re-Heat Treat- Studies, Surface Effects, and Re-Heat Treatment, by Jeremy V. Gluck and James W. ment, by Jeremy V. Gluck and James W. Freeman, November 1961. 103p. incl. figs. Freeman, November 1961. 103p. incl. fJs. tables and refs. (Project 7381; Task 73810) tables and refs. (Project 7381; Task 73810) (ASD TR 61-73, Pt. II) (Contract AF 33(616)- (ASD TR 61-73, Pt. II) (Contract AF 33(616)I64~62)' 662) Unclassified report Unclassified report The effect of creep-exposure on room temp- The effect of creep-exposure on room temperature mechanical properties of Rene' 41 erature mechanical properties of Rene' 41 was studied for temperatures of 1200-18000F was studied for temperatures of 1200-18000F and times up to 200 hours. Unstressed ex- and times up to 200 hours. Unstressed exposures were for as long as 2012 hours at posures were for as long as 2012 hours at 17000F. Thermally-induced structural 17000F. Thermally-induced structural changes reduced strength and ductility after changes reduced strength and ductility after exposures at 1400-18000F. Reduced yield exposures at 1400-18000F. Reduced yield UNCLASSIFIED UNCLASSIFIED over )( over ) UNCLASSIFIED UNCLASSIFIED strength was due to decreased volume fraction strength was due to decreased volume fraction of gamma prime and secondarily to an in- of gamma prime and secondarily to an increase in the particle size. Ductility was crease in the particle size, Ductility was reduced by formation of massive grain bound- reduced by formation of massive grain boundary carbides. Up to 15000F, creep caused ary carbides. Up to 15000F, creep caused strain hardening and Bauschinger effects. strain hardening and Bauschinger effects. Except for surface effects, damage was re- Except for surface effects, damage was restorable by re-heat treatment. Yield storable by re-heat treatment. Yield strength was restored by re-solution and re- strength was restored by re-solution and reaging to produce fine gamma prime. Complete I aging to produce fine gamma prime. Complete re-solution of carbides was required to re- re-solution of carbidus was required to restore ultimate strength and ductility. store ultimate strength and ductility, Microcracking was not observed. Creep in- Microcracking was not observed. Creep induced intergranular surface cracking at duced intergranular surface cracking at 1200-13000F which reduced ductility. Sur- 1200-13000F which reduced ductility. Surface effects for exposures above U1400F were face effects for exposures above 14000F were thermally induced. General principles were thermally induced. General principles were formulated for damage to properties of formulated for damage to properties of nickel-base alloys. UNCSFED nickel-base alloys. UNCA 4 —--- ---- -- -U NCAS S IF-ED I UNC -L-AS S IFIED

ASD TECHNICAL REPORT 61-73 PART II EFFECT OF CREEP-EXPOSURE ON MECHANICAL PROPERTIES OF RENE' 41 PART II: STRUCTURAL STUDIES, SURFACE EFFECTS, AND RE-HEAT TREATMENT JEREMY V. GLUCK JAMES W. FREEMAN THE UNIVERSITY OF MICHIGAN NOVEMBER 1961 DIRECTORATE OF MATERIALS AND PROCESSES CONTRACT No. AF 33(616)-6462 PROJECT No. 7381 AERONAUTICAL SYSTEMS DIVISION AIR FORCE SYSTEMS COMMAND UNITED STATES AIR FORCE WRIGHT-PATTERSON AIR FORCE BASE, OHIO 600 - January 1962 - 19-816 G 817

FOREWORD This report was prepared by The University of Michigan, Department of Chemical and Metallurgical Engineering under USAF Contract No. AF 33(616)-6462. This contract was conducted under Project No. 7381, "Materials Application", Task No. 73810, "Exploratory Design and Prototype Development". The work was administered under the direction of the Directorate of Materials and Processes, Deputy for Technology, Aeronautical Systems Division, with Messrs. W. H. Hill and D. M. Forney, Jr. acting as project engineers. This report covers work conducted between April 1, 1960 and May 31, 1961. The research is identified in the records of the University of Michigan as Project No. 02902. The invaluable assistance of Mr. P. D. Goodell in the structural studies is gratefully acknowledged.

ABSTRACT This report presents the results of a continued investigation of the influence of creep-exposure on the mechanical properties of Rene' 41 at room temperature beyond the results reported in ASD TR 61-73. The results show that thermally-induced structural changes reduced the strength and ductility for exposures at temperatures from 1400~ to 1800~F. Reduced yield strength was due mainly to a decrease in the measured volume fraction of gamma prime precipitate and secondarily to an increase in the gamma prime particle size. Exposures beyond 100-200 hours did not further reduce the yield strength due to the attainment of a near minimum volume fraction of gamma prime. For the conditions studied, gamma prime was sufficiently stable up to 1400~F so that yield strength was not affected. Ductility was reduced by thermally-induced carbide precipitation in the grain boundaries in the temperature range from 1400~ to 1800~F. Ductility reached a minimum after about 100 hours at 1700~ to 1800~F and was not apparently further decreased by creep. Creep also had little effect on the gamma prime reactions. Creep did introduce strain hardening which increased ultimate strength and decreased ductility and Bauschinger effects which raised the tensile yield strength for creep-exposures up to 1500~F. Creep also induced surface cracking which reduced ductility after exposure at 1200~ and 1300~F. Contact with alumel apparently accelerated the surface cracking. Where ductility was sufficiently reduced, the ultimate strength was also reduced. Thermally-induced surface reactions occurred for exposures above 1400~F. The maximum extent of such reduced ductility was attained rapidly and was independent of the exposure temperature. The removal of odd-size atom alloying elements from solution as a result of precipitation reactions was apparently responsible for a volume decrease during unstressed exposures above 1400~F. The damage effects, except for surface reactions, were not permanent. Strength could be restored by re-dissolving and re-precipitating gamma prime. Re-solution of carbides, particularly M6C by heating to 2150~F, was necessary to completely restore ultimate strength and ductility. Creep strength could also be restored. The absence of creep damage may have been due to the complete absence of microcracking during creep. The results are believed typical for nickel-base Ti-Al alloys except for the absence of the microcracking commonly occurring in many alloys of this type. A series of general principles was formulated for damage to mechanical properties of such alloys. The controlling influence of the volume fraction of gamma prime on the yield strength appears to be in agreement with theories of dispersion strengthening. Better verification of this than was possible during this investigation iii

ABSTRACT (continued) would be desirable. This is also true for the role of the types of carbides and the mechanisms of surface damage. PUBLICATION REVIEW This report has been reviewed and is approved. FOR THE COMMANDER: W. J. TRAPP Chief, Strength and Dynamics Branch Metals and Ceramics Laboratory Directorate of Materials and Processes iv

TABLE OF CONTENTS PAGE INTRODUCTION.......... 1 EXPERIMENTAL MATERIAL....... 2 TEST SPECIMENS........... 4 EQUIPMENT AND PROCEDURES...... 4 CREEP-EXPOSURE TESTS.............. 4 TENSILE TESTS.......... 5 IMPACT TESTS................... 6 DIMENSIONAL CHANGES.............. 6 STRUCTURAL EXAMINATION............. 7 Specimen Preparation........... 7 Etchant.................. 7 Optical Microscopy............ 7 Electron Microscopy............ 8 X-ray Diffraction............. 8 Lattice Parameter Determination.. 8 Minor Phase Identification....... 8 Electron Diffraction........... 9 RESULTS AND DISCUSSION EFFECT OF CREEP-EXPOSURE ON MECHANICAL PR.OPERTIES -- THERMALLY-INDUCED STRUCTURAL CHANGES...................... 11 Tensile Properties after Prolonged Exposure at 1600~, 1700~, and 1800~F............ 11 Volume Shrinkage During Exposure at 1600~, 1700~, and 1800~F................... 12 Impact Tests as a Measure of Creep Damage. 13 EFFECT OF RE-HEAT TREATMENT ON TENSILE PROPERTIES.......... 14 Re-Heat Treatment after Thermally-Induced Structural Changes............... 15 Re-Solution of Gamma Prime and M23C6 Carbides Only........ 15 Complete Re-Solution and Re-Heat Treatment. 15 Complete Re-Solution with Omission of 19500 and 1975 ~F Treatments........... 16 v

TABLE OF CONTENTS (continued) PAGE Complete Re-Solution Without Exposure.. 16 Re-Heat Treatment of Material Originally Heat Treated at 2050~ and 1650~F..... 16 Complete Re-Solution after Exposure to Creep at 1800~F. 17 Influence of Temperature and Time of Creep Exposure on Response to Heat Treatment... 17 Summary of Tensile Property Studies... 18 RE-HEAT TREATMENT AND CREEP- RUPTURE PROPERTIES.......... 19 R.e-Heat Treatment by Re-Solution of Gamma Prime and M23C6 Carbides Only........ 20 Complete Re-Solution and Re-Heat Treatment... 22 Influence of Surface Damage on Response to ReHeat Treatment 22 General Principles from Study of R.e-Heat Treatment on Creep-Rupture Properties.. 23 MICROSTRUCTURAL ASPECTS OF CREEP DAMAGE.. 25 Influence of Exposure and Re-Heat Treatment on Microstructure....... 26 Relation of the State of Gamma Prime after Exposure to Yield Strength..........27 Effect of Exposure on Matrix Lattice Parameter. 31 Relationship of Carbides to Properties.32 Additional Observations from Structural Studies.. 33 EFFECT ON TENSILE PROPERTIES OF SURFACE REACTIONS DURING EXPOSURE 34 Exposure Without Creep 34 Exposure With Creep 35 General Effects from As-Exposed Surfaces.... 35 Deep Cracking During Creep at 1200~ and 1300~F. 36 Examination of As-Exposed Surfaces.... 37 Cracking Characteristics.... 37 General Surface Alteration........ 38 Structural Characteristics of General Surface Alteration 39 GENERAL PRINCIPLES FOR CREEP-DAMAGE OF GAMMA PRIME-STRENGTHENED ALLOYS... 40 vi

TABLE OF CONTENTS (continued) PAGE CONCLUSIONS...................... 43 REFERENCES...................... 45 vii

LIST OF TABLES TABLE PAGE 1 Effect of Long-Time Aging on Tensile Properties of Rene' 41 47 2 Impact Test Data for Rene' 41 Tested at Room Temperature.......... 48 3 Establishment of Re-Heat Treatments for Restoration of Room Temperature Tensile Properties of Rene' 41. 49 4 Effect of Re-Heat Treatment on Tensile Properties of Rene' 41 After Initial Creep-Exposure.... 50 5 Effect of Re-Heat Treatment on Creep-Rupture Properties of Rene' 41 After Initial Creep Exposure. 51 6 Structural Parameters and Yield Strengths of Rene' 41 After Unstressed Exposure........... 52 7 Comparison Matrix Lattice Parameters and Dilatometer Shrinkage for Rene' 41 Aged Without Stress. 53 8 X-Ray Diffraction Data from Extraction Residues of Rene' 41 Specimens Used for Development of Re-Heat Treatments. 54 9 Effect of Re-Machining on Response of Room Temperature Tensile Properties of Rene' 41 to CreepExposure......... 55 10 Creep Exposure Test Data for Rene' 41....... 56 11 Cracking Depth Data for Creep-Exposure of Rene' 41 at 1200~ and 1300~F....... 57 12 General Surface Attack Data for Exposed Specimens of Rene' 41......... 58 13 X-Ray Diffraction Data from Extraction Residues of Dilatometer Aging Specimens........... 59 viii

LIST OF ILLUSTRATIONS FIGURE PAGE 1 Details of Test Specimens........... 60 2 Method of Attaching Thermocouple for Long-Time Dilatometric Aging Studies........... 61 3 Effect of Exposure Time in Unstressed Exposures on Room Temperature Tensile Properties of R.ene' 41 Alloy. ("R." Condition)............. 62 4 Shrinkage versus Aging Time for Dilatometer Exposures of Rene' 41....... 63 5 Comparison of Shrinkage Observed in Dilatometer Aging With Data Taken in Creep Units....... 64 6 Negative Creep Observed in Rene' 41 at 1200~F During Initial Survey Tests (R.ef. 1)...... 65 7 Effect of 10 Hours Unstressed Exposure on Smooth Bar Charpy Impact Strength of Rene' 41....... 66 8 Relative Efficiencies of Tensile and Smooth Bar Impact Strengths in Showing Damage to Rene' 41... 66 9 Effect of R.e-Heat Treatment on Tensile Properties of Rene' 41 After Creep-Exposure at 1400~F.. 67 10 Effect of Re-Heat Treatment on Tensile Properties of Rene' 41 After Creep-Exposure at 1600~ or 1700~F...................... 68 11 Effect of Re-Heat Treatment on Tensile Properties of Rene' 41 After Creep-Exposure at 1800~F.. 69 12 Effect of Creep-Exposure on Room Temperature Tensile Properties of Rene' 41 and Subsequent Ability to Recover Original Properties by Re-Heat Treating and Re-machining..... 70 ix

LIST OF ILLUSTRATIONS (continued) FIGURE PAGE 13 Effect of Re-Heat Treatment "R" on Rupture Life After Initial Creep-Exposure..... 71 14 Effect of Re-Heat Treatment to the "R" Condition and Re-machiningon Ratios of Actual Behavior to Theoretical Behavior of R.ene' 41 After a First Creep-Exposure of 10 or 100 Hours at 1400~, 1600~, or 1800~F......... 72 15 Effect of Re-Heat Treatment and Re-machiningon Creep of Rene' 41 After Initial Creep-Exposure for 6. 9-10 Hours........ 73 16 Effect of Re-Heat Treatment and Re-machiningon Creep of Rene' 41 After Initial Creep-Exposure for 100 Hours................... 74 17 Effect of Various R.e-Heat Treatments on Rupture Life of Rene' 41 at 1800~F......75 18 Effect of Type of Re-Heat Treatment (Followed By Be-machining) on Creep of Rene' 41 After Initial Creep-Exposure at 1800~F for 6.8-33 Hours.. 76 19 Effect of Type of Re-Heat Treatment (Followed By Re-machining) on Creep of Rene' 41 After Initial Creep-Exposure at 1800~F for 100 Hours.. 77 20 Effect of Re-Heat Treatment and Re-machiningon Creep of Rene' 41 After Initial Creep-Exposure at 1600~F................... 78 21 As-Heat Treated "R." Condition....... 79 22 Specimen R-104 Exposed Without Stress 100 Hours at 1800F..................... 79 23 Specimen R.-105 Re-Heat Treated After 100 Hours at 1800~F (Includes One Hour Re-Solution at 2150~F). 79 x

LIST OF ILLUSTRATIONS (continued) FIGURE PAGE 24 Specimen R.-lll R.e-Heat Treated After 100 Hours at 1800~F (Includes Two Hour Re-Solution at 2150~F) 79 25 Optical Micrograph of Spec. R-150 Creep-Exposure 10 Hours at 1400~F to 3. 52 Percent Deformation, Then Re-Heat Treated to Condition "R.".. 80 26 Electron Micrograph of Spec. R-150..... 80 27 Electron Micrograph of Spec. R-172 Creep-Exposure 100 Hours at 1400~F to 19. 0 Percent Deformation, Then R.e-Heat Treated to Condition "R.".. 80 28 Electron Micrograph of Spec. R.-148 Creep-Exposure 10 Hours at 1600~F to 6. 20 Percent Deformation, Then Re-Heat Treated to Condition "R.... 80 29 Electron Micrograph of Spec. R-185 Creep-Exposure 100 Hours at 1600~F to 7. 38 Percent Deformation, Then Re-Heat Treated to Condition "R.".. 81 30 Electron Micrograph of Spec. R-151 Creep-Exposure 10 Hours at 1700~F to 3. 57 Percent Deformation, Then Re-Heat Treated to Condition "R"... 81 31 Electron Micrograph of Spec. R-155 Creep-Exposure 10 Hours at 1800~F to 5. 34 Percent Deformation, Then Re-Heat Treated to Condition "R.... 81 32 Optical Micrograph of Spec. R-197 Creep-Exposure 100 Hours at 1800~F to 8. 06 Percent Deformation, Then Re-Heat Treated to Condition "R.".. 82 33 Electron Micrograph of Spec. R-197... 82 34 Optical Micrograph of Spec. R-124 Creep-Exposure 100 Hours at 1800~F to 3.46 Percent Deformation, Then Re-Heat Treated to Condition A".... 82 xi

LIST OF ILLUSTRATIONS (continued) FIGURE PAGE 35 Electron Micrograph of Spec. R-124.. 82 36 Effect of Aging Time and Temperature on Average Gamma Prime Particle Size in Rene' 41.. 83 37 Effect of Exposure Time on Volume Fraction of Gamma Prime in Matrix of Rene' 41.. 84 38 Effect of Gamma Prime Particle Size on Yield Strength of Rene' 41.. 85 39 Correlation of Meiklejohn-Skoda Volume Fraction Parameter With Room Temperature Yield Strength of Rene' 41 After Unstressed Exposure... 86 40 Collodion Replica of Spec. R-32........ 87 41 Carbon Extraction Replica of Spec. R-32.... 87 42 Carbon Extraction Replica of Spec. R-32.... 87 43 Collodion Replica of Spec. R-104. Spec. R-32 and Spec. R-104 Exposed Without Stress 100 Hours at 1800~F..........87 44 Collodion Replica of Spec. R-104........ 88 45 Carbon Extraction Replica of Spec. R-104.... 88 46 Carbon Extraction Replica of Spec. R-104.... 88 47 Carbon Extraction Replica of Spec. R-104 (Diffraction obtained from outlined area) Spec. R-104 Exposed Without Stress 100 Hours at 1800~F... 88 48 Carbon Extraction Replica of Spec. R-104 (Diffraction obtained from outlined area).... 89 49 Electron Diffraction Photo of Selected Area of Figure 47 89 xii

LIST OF ILLUSTRATIONS (continued) FIGURE PAGE 50 Electron Micrograph of As-Treated Condition "R" (1950~F Solution)............. 89 51 Electron Micrograph of Spec. R.-124 Re-Heat Treated to Condition "A" (2150~F Solution) After 100 Hours Creep-Exposure at 1800~F..... 89 52 Effect of Re-machining on Curves of Room Temperature Tensile Properties Versus Exposure Temperature of R.ene' 41 Exposed Without Stress for 10 or 100 Hours................. 90 53a Effect of Exposure Temperature and Time on Difference in Room Temperature Tensile Properties Between As-Exposed and Re-machined Specimens of Rene' 41 After Unstressed Exposure 91 53b Effect of Exposure Temperature and Amount of Prior Creep on Difference in Room Temperature Tensile Properties Between As-Exposed and Re-machinedSpecimens of R.ene' 41 After 10 Hours Exposure........ 91 54 Effect of Prior Creep, Re-machining, and Thermocouple Practice on Room Temperature Tensile Properties and Cracking Tendencies of Rene' 41 Exposed at 1200~ and 1300~F. 92 55 Effect of Prior Creep and Re-machining on Room Temperature Tensile Properties of Rene' 41 Exposed at 1400~, 1600~, and 1800~F... 93 56 Examples of Deep Cracks in Rene' 41 Exposed to Creep at 1200~ or 1300~F (Low Ductility and Oxidation on Fracture Surface for Subsequent Tensile Test at Room Temperature).. 94 57 Examples of Creep Curves of Rene' 41 Specimens Exhibiting Deep Cracking After 10 Hours CreepExposure at 1200~ or 1300~F......... 95 xiii

LIST OF ILLUSTRATIONS (continued) FIGURE PAGE 58 Spec. No. R-90 (Section Showing Surface) Exposed Without Stress 100 Hours at 1000~F.. 96 59 Spec. No. R-99 (Section Showing Surface Attack) Creep-Exposure at 1300~F for 100 Hours to 0. 63 Percent Deformation.............96 60 Spec. No. R-95 (Section Showing Surface Attack and Isolated Crack) Creep-Exposure at 1300~F for 10 Hours to 2. 71 Percent Deformation.. 96 61 Spec. No. R-93 (Section Showing Surface Attack) Exposed Without Stress 100 Hours at 1400~F.. 97 62 Spec. No. R-92 (Section Showing Surface Attack) Exposed Without Stress 100 Hours at 1600~F.. 97 63 Spec. No. R-91 (Section Showing Surface Attack) Exposed Without Stress 100 Hours at 1800~F.. 97 64 Longitudinal Section of Spec. D-2 Exposed Without Stress 474 Hours at 1700~F.... 98 65 Electron Micrograph of Spec. D-2...... 98 66 Longitudinal Section of Spec. D-5 Exposed Without Stress 2012 Hours at 1700~F..... 98 67 Electron Micrograph of Spec. D-5...... 98 68 Effect of Exposure Temperature and Time on Depth of Visible Surface Attack in Rene' 41.. 99 69 Effect of Visible Surface Attack on Room Temperature Tensile Ductility of Rene' 41 Exposed Without Stress at 1300~ - 1800~F 100 70 Section of Spec. D-2 Showing Attacked Surface Layers (Arrow shows solution of grain boundary phase in advance of general attack)........... 101 xiv

LIST OF ILLUSTRATIONS (continued) FIGURE PAGE 71 Electron Micrograph of "Depleted" Layer of Spec. D-2 Exposed 474 Hours at 1700~F in Dilatometer Furnace.......... 101 72 Section of Spec. D-5 Showing Attacked Surface Layer (Exposed 1150 Hours at 1800~F in Dilatometer).......... 102 73 Effect of Surface Attack on Matrix Lattice Parameter of Spec. D-2............... 103 74 Lattice Parameter Data for Specimen D-2 Compared to Data of Beattie and VerSnyder for Effect of Molybdenum Content on Lattice Parameter of NickelBase Alloys.................. 103 xv

INTRODUCTION An investigation has been conducted at the University of Michigan under the sponsorship of the Materials Central, ASD, United States Air Force, under Contract AF 33(616)-6462 to obtain information and develop general principles for the influence of creep-exposure on the normal temperature mechanical properties of alloys. Changes in mechanical properties are termed "creep damage" and may be defined as any degradation in properties following exposure to stress and/or temperature under which creep can occur. This report presents data on this subject for Rene' 41, a nickel-base Al+Ti-hardened superalloy. The results of research covering the effects of 10 to 200 hours creep-exposure over the temperature range from 1200~ to 1800~F were presented previously (Ref. 1) for conditions where surface damage effects were excluded, as far as possible, from the experimental variables. The initial results indicated that both thermally-induced structural changes and creep-strain caused alterations of room temperature mechanical properties. In addition, there was an indication that surface effects could occur that were almost as large as the thermally-induced changes, and furthermore, probably operated at lower temperatures. The major thermally-induced structural change identified was the agglomeration of dispersed gamma prime particles in the matrix. The metal carbides M6C and M23C6 increased during creep-exposure and there were indications that the carbides were an important factor controlling ductility. Strain hardening and the Bauschinger effect contributed to increased strength and decreased ductility. The research described in this report was designed to obtain additional mechanical property and microstructural information to supplement the results of Reference 1 and to extend the investigation by studying the influence of surface reactions during exposure and the feasibility of restoring properties by re-heat treatment. Earlier research (Refs. 2-5) was confined to studies of the effects of specific exposure conditions on a number of established alloys used in sheet form in aircraft structures. The present investigation deviated from this practice by deliberately seeking damage from exposure conditions without restriction to engineering limits commonly used to avoid damage. In order Manuscript released for publication June 27, 1961 as an ASD Technical Report. 1

to induce changes in Rene' 41, extremes of both stress and temperature were utilized to outline exposure principles causing damage. The material was particularly outstanding in its lack of susceptibility to internal microcracking after substantial amounts of creep at high temperatures. The studies of the influence of creep-exposure on room temperature properties were based mainly on tensile testing. Some compression testing and impact testing was also conducted. Surface effects were eliminated in the initial studies by re-machining the specimens following creep-exposure. Following mechanical property determinations, extensive microstructural studies were conducted in order to define the mechanisms producing the damage. Use was made of optical and electron microscopy and x-ray and electron diffraction. The selection of Rene' 41 as an experimental material was made to provide specific data for the alloy and to develop general principles for the numerous nickel-base alloys of the same metallurgical type that might be used in aircraft structures. EXPERIMENTAL MATERIAL The primary requirement for the experimental material was that it be a complex high-strength heat resistant alloy with considerable future applicability by the Air Force. Preliminary experiments, discussed in detail in Reference 1, led to the selection of the nickel-base alloy Rene' 41 as a promising representative of this class of material. Approximately 107 pounds of Rene' 41 alloy were procured from the Metallurgical Products Department, General Electric Company, in the form of 0. 516-0. 520-inch diameter centerless ground bar stock from vacuum induction-melted Heat R-134. The material was reported to have the following chemical composition: Element Weight Percent Nickel 55. 32 (by difference) Chromium 19. 27 Cobalt 11.06 Molybdenum 9. 06 Titanium 3. 28 Aluminum 1. 44 Carbon 0. 12 2

Chemical Composition for Rene' 41 (Continued) Element Weight Percent Boron 0. 0040 Iron <0. 30 Sulfur 0. 006 Manganese 0.07 Silicon 0. 07 The producer reported the grain size to be ASTM 3-6 and the Brinell Hardness to be 241-269. The material was shipped in the mill annealed condition (1975~F solution treatment plus water quench). Most of the material was tested in the heat treatment condition recommended by General Electric for applications limited in service by tensile properties (R.ef. 6). This treatment, coded "R" for the purposes of the investigation, consisted of the following: 1) Solution treatment: 1950~F - 1/2 hour plus air cool 2) Aging treatment: 1400~F - 16 hours plus air cool The Vickers Hardness after treatment was 367 (equivalent to a Brinell Hardness of 347). A minor amount of experimental work was performed on specimens heat treated according to a specification by General Electric for rupturelimited applications (Ref. 6). This treatment was coded tR2?" and consisted of the following: 1) Solution treatment: 2050~F - 1/2 hour plus air cool 2) Aging treatment: 1650~F - 4 hours plus air cool This treatment produced a Vickers Hardness of 348 and was shown in Reference 1 to produce substantially the same short-time creep-rupture properties as the "R." treatment and a similar response to creep-exposure. For the original heat treatments, batches of 24 specimen blanks, 4 inches long by 0. 516-inch diameter, were heat treated in 4 bundles of 6 blanks each. Re-heat treatments were generally conducted as convenient -- in some cases, groups of up to 6 specimens were treated at one time. The details of the re-heat treatments will be discussed more fully in the section on Results. 3

TEST SPECIMENS Details of the test specimens employed in this investigation are shown in Figure 1. The original creep-exposures were generally conducted on specimens having a 0. 350-inch diameter gage section with a reduced section approximately 2 inches long. For high stress tests, gage section diameters of 0. 300-inches were occasionally used in order to reduce the absolute load on the creep testing machine. The specimens for subsequent mechanical property tests were designed to be machined from the gage sections of the creep-exposure specimens. Where surface damage was to be eliminated as a variable, approximately 0. 025-inches were machined from the gage section diameter. Following rough turning on a lathe, all gage sections were draw-filed and then hand polished with crocus cloth. The impact specimen employed for smooth bar tests was the ASTM Type-W sub-sized bar. The dimensions, 0. 197-inches square by 2. 16 inches long, were such that it could be conveniently milled from the reduced section of a creep specimen. Exposure tests without stress were also conducted on cylindrical specimens 4. 0-inches long by 0. 4-inches in diameter in a dilatometer so that volume changes could be measured. Care was taken that the specimen ends were flat and parallel. EQUIPMENT AND PROCEDURES CREEP-EXPOSURE TESTS The creep-exposure tests were conducted in individual University of Michigan creep-testing machines. In these units, the stress is applied through a third class lever system having a lever-arm ratio of about 10 to 1. The specimens were gripped by threaded holders fitting into a universal joint system that insured uniaxial loading. Heating was accomplished by a wire-wound resistance furnace fitting over the entire specimen holder assembly. Strain measurements were made with the modified Martens extensometer system, an optical lever-arm system which permits the detection of specimen elongations of approximately 10 millionths of an inch. Because this system required the attachment of extensometer bars to collars threaded to the specimen shoulders, it was necessary to correct the observed deformations for the diminished contribution of the fillets and shoulders. A detailed description of the extensometer system and the method of making 4

"effective gage length" calculations was given previously (R.efs. 1 and 3) and will not be repeated here. Temperature measurements were generally made from three Type-K thermocouples, one at the center of the reduced section and at either end. All thermocouples were shielded from direct furnace radiation. Prior to starting a test, the furnace was heated to within 50~F of the desired temperature. The specimen was then placed in the hot furnace and brought up to the test temperature and distribution in a standard period of four hours. ASTM Recommended Practices were followed in controlling the test temperature and distribution. For several tests where exposure was desired without thermocouple contact on the gage section, temperature measurement was made from couples wired to the specimen shoulders. The validity of this procedure was checked on a specimen with five couples attached -- the three on the gage section and one at either shoulder. The shoulder-toshoulder temperature deviation was less than 5~F Strain measurements were made as each weight was applied during loading and then periodically throughout the tests. At the end of the exposure period, a final reading was made and the power to the furnace turned off. The specimen then cooled under load in order to minimize the effects of creep recovery. In the tests exposed without stress, the same-procedure was followed in order to make the total time at temperature equivalent to the case of a stressed exposure. The Martens extensometers were also used in the unstressed exposures as an expedient to obtain information on dimensional changes due to structural reactions. For the 0. 350-inch diameter specimens, the weight of the necessary holder and extensometer system produced a specimen stress of 54. 5 pounds per square inch. The data obtained with this system were later checked by conventional dilatometer tests. Where the creep tests were allowed to run until failure, an automatic timer, accurate to one-tenth of an hour, was actuated by the fall of the specimen holder to measure the rupture time. TENSILE TESTS Short-time tensile tests were conducted at room temperature in a hydraulic testing machine. Elongation was measured with a microformer strain gage and recorded automatically in the form of a load versus strain curve. A strain pacer was used to insure a strain rate of 0. 005 inches per inch per minute. The data determined in the tensile tests were the ultimate tensile 5

strength, the 0. 2-percent offset yield strength, the elongation, reduction of area, and the modulus. IMPACT TESTS Both Izod and Charpy smooth bar impact tests were utilized in this investigation. In the Izod test, the specimen is supported vertically at one end as a cantilever, while in the Charpy test the specimen is supported horizontally at both ends as a simple beam. A special holding fixture was used to accommodate the sub-sized specimen. In all tests, the impact machine pendulum was set to produce a striking energy of 120 foot-pounds Prior to conducting a test, the scale of the machine was zeroed by allowing the pendulum to swing through its cycle with no specimen in place. DIMENSIONAL CHANGES The dimensional change data obtained using the optical extensometer system during the unstressed exposures were checked by using more conventional dilatometric techniques. A quartz tube dilatometer assembly conforming to ASTM Recommended Practices was available for these determinations. A wire-wound resistance furnace was used to heat the assembly. The test specimen employed was a cylinder 4. 0-inches long by 0. 4-inches in diameter. Care was taken that the ends of the specimen were flat, parallel, and perpendicular to the specimen axis. Dimensional changes were measured with a mechanical dial gage graduated in 10, 000ths of an inch. Since the aim of these tests was to obtain information on dimensional changes with time at temperature and not the thermal coefficient of expansion, minor modifications of conventional procedure were employed. To reduce the effects of heating time, the tests were hot started, that is, the specimens were placed in the hot unit which had been previously brought to about 10~F below the desired aging temperature. The specimen was then "nursed" to the proper temperature as quickly as possible, generally 40 minutes or less, and the dial gage readings were started. From two to four readings per day were then made throughout the exposure period. Temperature measurement was made from a thermocouple attached to the center of the specimen. It was not possible to spot weld a thermocouple to the specimen because the 28-gage wire necessary for a couple that could be spot welded would not stand up during the long aging times contemplated. The standard dilatometer rod was modified to permit the use of 18-gage thermocouples as shown in Figure 2. This was accomplished by using a hollow quartz tube for the push rod. The heavy gage thermocouple wires were covered with ceramic insulators and fed through the push rod. At the bottom end, the bare wires were brought out around the specimen, 6

fusion welded to form a bead, and then fastened to the specimen with asbestos cord. The top of the thermocouple was brought out around a special cap. A thermocouple that had been attached in this manner to a specimen during aging for 2012 hours at 1700~F was subsequently checked against a virgin couple at 1700~F and found to read only 5~F hot. STRUCTURAL EXAMINATION The techniques used in this investigation for structural examination included optical microscopy, electron microscopy, x-ray diffraction of powder samples and extracted residues, and electron diffraction of extraction replicas. Specimen Preparation Specimens for microscopic examination were sectioned longitudinally along the center line with a water-cooled cut-off wheel and then mounted in Bakelite. The mounted specimens were wet ground on a rotating lap through a series of silicon carbide papers finishing at 600-mesh grit. Final polishing was accomplished on a cloth-covered rotating lap using fine diamond compound and then on a vibratory polisher in an aqueous media of Linde I"B" polishing compound. The samples were cleaned in an ultrasonic cleaner with a detergent solution. Etchant The specimens for optical and electron microscopy were etched electrolytically in "G" etch, an etchant developed by Bigelow, Amy, and Brockway (Ref. 7). Etching was conducted at 6 volts and a current density of approximately 0. 8 amperes per square inch for a period of 5-7 seconds. The composition of the etchant follows: "G" Etch H3P04 (85%) 12 parts H2S04 (96%) 47 parts HN03 (70%) 41 parts Optical Microscopy Conventional methods were employed for general optical examination. Crack depths and oxide penetrations were determined with a 10x filar micrometer eyepiece at a system magnification of 600x. 7

Electron Microscopy Electron microscopy was carried out in an R.CA EML Electron Microscope. For examination in the microscope, collodion replicas mounted on nickel grids were prepared from the surface of the etched specimens. The replicas were shadowed with palladium to increase the contrast and reveal surface contours. Polystyrene latex spheres of either 2580 or 3400 A diameter were placed on the replicas prior to shadowing to indicate the angle and direction of shadowing and to provide an internal standard for the measurement of magnification. The micrographs reproduced in this report are direct prints from the original negatives; consequently, the polystyrene spheres appear black and the shadows appear white. Since the spheres are raised from the surface of the replica, a particle casting a shadow opposite to that of the latex spheres is in relief on the metal specimen; conversely, areas casting shadows in the same direction as those cast by the latex spheres are depressions in the surface of the metal specimen and correspond to material that was attacked or eaten out by the etchant. X-ray Diffraction X-ray diffraction analysis was used both for the identification of minor phases and the determination of matrix lattice parameters. Lattice Parameter Determination. Determinations of the matrix lattice parameters were made from Debye-Scherer diffraction photographs of filings. The filings were made with a very fine file, sifted with a magnet to remove any extraneous iron particles that might have come from the file, and then sifted through a 200-mesh screen. The resultant fine powder was rolled into a thin wire using a binder of Duco cement. Exposure was made to nickel-filtered copper radiation for 4 hours in a 114. 6 mm diameter Debye camera. An optical comparator was used to determine diffraction line positions. Lattice parameters were determined by the extrapolation method of Nelson and Riley (Ref. 8). Minor Phase Identification. The identification of minor phases was made from x-ray diffraction analysis of residues extracted either from solid samples or from filings by immersion in a bromine-alcohol solution. Filings were made for the identification of the constituents in various surface layers of an aged dilatometer sample. Layers approximately onethousandth of an inch thick were removed from the sample with a fine file as it was rotated in a lathe. After initial preparation, the samples for extraction were soaked in pure bromine liquid, cleaned with distilled water, and completely dried. 8

Each specimen was then placed in an individual centrifuge tube. Approximately 25 ml. of a 10:1 mixture of anhydrous methyl alcohol and bromine was added. Boiling took place in about 10 minutes, after which the tube was cooled under tap water. The reaction was allowed to continue until a sufficient quantity of extract had been obtained, after which the specimen was removed and rinsed with methyl alcohol. The tube was centrifuged and the extract washed repeatedly with alcohol until the supernatant liquid was clear. The extract was scraped onto a filter paper, dried, and formed into a thin wire using a Duco cement binder. X-ray exposures were conducted either in a 57. 3 or a 144. 6 mm diameter Debye camera using nickelfiltered copper radiation for periods of four hours. The line positions were determined on an optical comparator and dit' values were calculated. The patterns were then analyzed by comparison with standard patterns available from the literature and, in particular, the ASTM Powder Data File. Electron Diffraction Electron diffraction studies were carried out in a 50 kilo-volt RCA EML Electron Microscope on extraction replicas mounted on nickel grids. Although this process has the advantage of allowing study of phases in situ, it is a time-consuming, rather delicate operation, particularly sensitive to preparation techniques. Results once obtained, must be analyzed carefully. Considerable experimentation was required before usable replicas could be prepared, with best results finally being obtained from bromineextracted carbon replicas. Replicating media discarded due to inability to find a satisfactory extracting agent that would not also destroy the replica included collodion and Fax-film. A vinyl chloride polymer appeared to have promise in standing up to a number of extracting reagents, however, the polymerization often continued to the point of cracking. The metallographic specimens selected for extraction replication were repolished and etched with "GI' etch in order to bring the minor phases well into relief. On some samples, the etching conditions were varied to also etch out the gamma prime phase. A carbon film was then vapor-deposited on the surface. This acted not only as a replica of the surface but was also the medium for holding the extracted particles in place. A 3/16-inch grid was then scribed in the carbon film with a razor blade to permit the penetration of the extracting agent to the specimen surface and to facilitate handling of the individual replicas. The mounted specimen was then immersed in a 10:1 bromine-absolute methyl alcohol solution for 4-6 hours. Following extraction, individual grid squares were floated off the surface by immersing the specimen in distilled water. The squares were washed and lifted out of the solution on nickel grids. 9

Each grid was dried and a thin film of aluminum was vapor-deposited on the replica to act as an internal standard for diffraction analysis. The completed grids were handled in the microscope in the same manner as conventional collodion replicas. An electron micrograph was taken to identify each area from which diffraction was obtained. The relatively low accelerating potential of the microscope made it difficult to obtain a diffraction pattern from anything other than thin films or particles. The completed patterns were analyzed by indexing the diffraction spots by a trial and error procedure analogous to the study of a frontreflection Laue x-ray pattern. The lattice constant of the diffracting phase was determined by the relationship of its pattern to the known pattern produced by the aluminum standard. 10

RESULTS AND DISCUSSION The previous report (Ref. 1) demonstrated that creep damage to mechanical properties at room temperature occurred in Rene' 41 by: (1) Residual mechanical effects from creep which caused Bauschinger effects and strain hardening (2) Thermally-induced structural changes (3) Surface reactions in air. This report extends the data for the second and third effects and presents results of investigations of the causes for the observed changes in properties as a basis for developing general principles for creep damage in gamma prime - strengthened alloys. EFFECT OF CREEP-EXPOSURE ON MECHANICAL PROPERTIES -- THERMALLY -INDUCED STRUCTURAL CHANGES A study was carried out of the effect of prolonging exposure time on tensile properties. A short study of the changes in impact properties for unnotched specimens was conducted. Volume shrinkage during exposure was studied. The effects studied were confined to thermally-induced structural changes by remachining specimen surfaces after exposure. Tensile Properties after Prolonged Exposure at 1600~, 1700~, and 1800~F Prolonging exposures at 1700~ and 1800~F without stress did not appreciably reduce tensile properties at room temperature (Table 1 and Fig. 3) from those resulting from the 100-200 hour exposures reported in Reference 1. The data covered exposures as long as 2012 hours at 1700~F and 1700 hours at 1800~F. Short-time exposures at 1600~F did not reduce properties as much as they did at 1700~ and 1800~F. An exposure of 401 hours at 1600~F reduced the properties to those after exposure for 100 to 200 hours at 1700~ and 1800~F. Longer exposures at 1600~F were not available to establish whether or not the properties would also level off at the same value as for exposure at 1700~ and 1800~F. The approximate levels of tensile properties at room temperature for exposures longer than 100 to 200 hours at 1700~ and 1800~F and for 400 hours at 1600~F in comparison to original material were as follows: 11

Property Exposed Material Original Material Ultimate strength (psi) 140, 000 190, 000 0. 2% offset yield strength (psi) 95, 000 130, 000 Elongation (%) 5 20 Reduction of area (%) 6 28 Values somewhat above and below these approximate levels after exposures were obtained in the actual tests. All exposures longer than 200 hours were carried out in a dilatometer so that volume changes could be measured. Those up to 200 hours were carried out in creep units under a dead load of 54. 5 psi. Volume Shrinkage During Exposure at 1600~, 1700~, and 1800~F The prolonged exposures without stress were carried out in a thermal expansion unit to measure volume changes accompanying structural changes. The data obtained (Fig. 4) indicated continued shrinkage at least up to 1150 hours at 1800~F and up to 2012 hours at 1700~F. Exposure for 401 hours at 1600~F also resulted in shrinkage. The amount of shrinkage was larger for a given time of exposure the higher the exposure temperature. The rate of shrinkage decreased with time but continued for the full period of exposure. Exposure for 1000 hours at 1800~F resulted in a shrinkage of 0. 004 inch and 0. 0028 inch at 1700~F. A similar value for 1600~F was not established but would apparently have been less than at 1700~F. These data indicate that there was a real and substantial shrinkage during exposure at 1600~ to 1800~F. Due to the stresses which can be induced by shrinkage, the characteristics of the shrinkage are in themselves important. It will be noted, however, that shrinkage continued after there was little further change in tensile properties at room temperature with continued exposure (Fig. 3). It is, therefore, uncertain as to what degree the structural changes causing shrinkage were responsible for the changes in mechanical properties. The shrinkage offsets and reduces the amount of creep measured during stressed exposure. Figure 5 compares creep curves with the shrinkage during unstressed exposure for the first 100 to 200 hours of exposure. This comparison indicates the amount the measured creep was reduced by the shrinkage. It does not, however, show whether or not stress accelerated shrinkage. The data of Figure 5 indicated that the unstressed tests reported in Reference 1 resulted in less shrinkage that was measured by the dilatometer. The testing procedure used for the previous tests was then 12

checked. These exposures were carried out in creep units with all adapters in place so that the extensometer system would operate. The stress imposed by the adapter system was 54. 5 psi. This was not sufficient stress to account for the difference from the amounts measured in the dilatometer. It is known that the extensometer method of measurement is subject to errors under loads too low to eliminate the misalignment associated with joints. It is, therefore, considered that the initial data correctly indicated shrinkage but that due to the conditions of measurement, the absolute amount was slightly low. In fact, one of the reasons for using the dilatometer was the question regarding the prior method of measuring the shrinkage. This allowed verification of the shrinkage by another method. The unstressed exposures at 1200~ and 1400~F in the creep units indicated little change in length in 100 to 200 hours. However, creep curves for stressed exposure at 75, 000 psi at 1200~F did show a small amount of shrinkage, exhibited as "negative creep" (Fig. 6). The phenomenon of "negative creep" has also been reported in the literature for a number of other alloy systems (Refs. 9, 10). Presumably, the shrinkage offsets some of the normal creep at higher stresses so that the amount of creep measured, such as for the curve at 114, 000 psi in Figure 6, was less than actually occurred. It is assumed that the normal creep at 75, 000 psi was extremely small and the indicated shrinkage was about the total amount which occurred. The indicated shrinkage probably represents the extensometer functioning correctly under load, whereas it did not in the unstressed exposures. It is possible, however, that the stress did accelerate shrinkage. Impact Tests as a Measure of Creep Damage Impact tests are frequently a most sensitive method of determining that properties are being changed as a result of exposure to creep. This method was abandoned in the initial studies (Ref. 1) when it was discovered that the subsize notched specimens had only about 3 foot-pounds of energy absorption in the as-treated condition. Subsequent discussion of this test with a representative of the Materials Central suggested that impact tests on unnotched specimens might provide useful results. The usefulness of this test was studied with the data obtained being given in Table 2 along with the original data for notched specimens from Reference 1. The unnotched specimens bent in the Izod tests at 30 footpounds and fractured at about 29 foot-pounds in Charpy tests. These values were sufficiently large so that reductions in properties due to exposure to creep might be useful for measuring damage. 13

Charpy tests on smooth specimens after unstressed exposure (Table 2 and Fig. 7) gave the following results: (1) Exposure for 10 hours at 1600~F did not significantly change impact strength. When the time was increased to 401 hours, there was a significant decrease. (2) Exposure for 10 hours at 1700~ or 1800~F did reduce impact strength significantly. (3) The scatter between individual tests could be quite large. For this reason, replicate testing would be necessary. Comparison of the information obtained from tensile and impact tests (Fig. 8) using average percentages of as-treated strength indicated a more pronounced decrease for impact properties. The impact strength of the smooth specimens was approaching low values so rapidly that with slightly more exposure all values would be very low and not show any change. As previously shown (Fig. 3), tensile properties continued to decrease with longer exposure. It was concluded that the impact tests were not showing changes where tensile tests did not show any change. The replicate exposures necessary were also undesirable. For these reasons, it was decided to discontinue impact tests. EFFECT OF RE-HEAT TREATMENT ON TENSILE PROPERTIES The study of the effects of re-heat treatments after creep-exposure was undertaken for the information it would provide regarding the factors causing creep damage to tensile properties at room temperature. The experiments had the following objectives: (1) To determine if the damage was permanent or if the properties could be restored by re-heat treatment. If any part of the damage were permanent, it would indicate that damage by internal cracking could have been a factor in addition to microstructural effects. If properties could be restored, it would indicate that microstructural changes alone were responsible. (2) To obtain information on the damage mechanism by varying the conditions of re-heat treatment. Beattie (Ref. 11) had determined for Rene' 41 that gamma prime dissolves between 1900~ to 1950~F, M23C6 carbides dissolve between 1700~ and 1800~F, and 2150~F is required to dissolve M6C carbides. For the present studies, the assumption was made that re-solution plus re-aging would result in nearly the same structure as the original material. 14

In all experiments, the specimens were re-machined after heat treatment. Surface effects from the prior creep-exposure as well as those from the re-heat treatment were thus eliminated. The tensile properties were calculated on the basis of the specimen dimensions after re-machining, The results will show that re-heat treatment was very effective in restoring properties at room temperature. It must be clearly recognized, however, that restoration would not be nearly as effective if the damaged surface is not removed. Re-Heat Treatment after Thermally-Induced Structural Changes The first step in the studies of re-heat treating was to evaluate the permanency of thermally-induced structural changes. The data are summarized in Table 3. Severe reduction in room temperature properties was induced by unstressed exposures at 100 hours at 1800~F (Table 3, Spec. R.-32 and Spec. R-104 indicate the effect of this exposure). In all the re-heat treatments studied, complete re-solution of the gamma prime was then induced. By varying the temperature of re-solution, the effect of leaving the M6C undissolved or dissolved was studied. Re-Solution of Gamma Prime and M23C6 Carbides Only: Re-heat treatment by reproducing the original treatment of one-half hour at 1950~F, air cooling and aging at 1400~F for 16 hours (coded "R" in this investigation) would be expected to re-dissolve and re-precipitate gamma prime as a fine dispersion. Heating at 1950~F would also be expected to dissolve at least part of any M23C6 carbides present. The treatment brought the yield strength up to nearly the original level (Table 3, Spec. R.-136). However, the ultimate strength was only partially restored and the ductility was not restored at all. This suggests that the solution of gamma prime was effective in restoring yield strength, but incomplete solution of carbides, particularly M6C, restricted ductility and thereby ultimate strength. Complete Re-Solution and R.e-Heat Treatment: After 100 hours exposure at 1800~F, Specimens R-105 and R-110 were heated to 2150~F and air cooled after 1 hour in one case and 2 hours in another. They were then re-heated at 1975~F for 1 hour and water quenched to simulate the original mill anneal. The original "R" heat treatment was then duplicated by heating to 1950~F for one-half hour, air cooling and aging at 1400~F for 16 hours. Both treatments nearly restored the original properties with the two-hour treatment at 2150~F (Spec. R-110) possibly slightly more effective. Taking into 15

account the probable differences in structure between the original and reheat treatments, it can be concluded that the effects of the thermallyinduced structural changes were eliminated and the original properties completely restored. Complete Re-Solution with Omission of 1950~ and 1975~F Treatments: After 100 hours exposure at 1800~F, re-treatment at 2150~F for 2 hours, plus simulation of the mill anneal at 1975~F and aging at 1400~F for 16 hours resulted in essentially the same properties (Table 3, Spec. R-lll and Spec. R-201) as the complete treatment. Ductility may have been sufficiently high to represent a significant increase. This treatment was coded "A". Omission of both the 1950~ and 1975~F treatments resulted in essentially the same properties (Spec. R-135). Evidently, the intermediate treatments at 1975~ and 1950~F had very little effect on tensile properties at room temperature. If these treatments did cause some re-precipitationof M6C while at 1975~ and 1950 F, it did not have much effect. Complete Re-Solution Without Exposure: Because the treatment at 2150~F was so effective after the 100-hour exposure at 1800~F, it seemed desirable to determine what the effect of this treatment would be if no exposure was involved. Accordingly, Specimen R-141 in the as-treated "R." condition was re-heated for 2 hours at 2150~F, air cooled, re-heated to 1975~F, water quenched and aged at 1400~F for 16 hours. Ultimate and yield strengths were lower than for the original material or the exposed specimens given the same re-treatment, while ductility was higher. While only one test was involved, it seems evident that the prior exposure at 1800~F for 100 hours resulted in some effect that slightly raised tensile strength and slightly reduced ductility in relation to unexposed material with the same treatment. Re-Heat Treatment of Material Originally Heat Treated at 2050~ and 1650 F: The heat treatment of one-half hour at 2050 ~F, air cool plus aging for 4 hours at 1650~F (code "R2") which had been recommended for rupture-limited applications (Ref. 6) was applied to several specimens. These were then exposed at 1800~F for 100 hours to induce marked reduction of tensile properties at room temperature (Table 3, Spec. R.2-8). Re-heat treatment with the original conditions of heat treatment restored yield strength but not ultimate strength and ductility (Table 3, Spec. R2-13). Evidently, treatment at 2050~F was not sufficiently high to dissolve enough carbides to restore properties. 16

R.e-heat treatment at 2150~F with or without the treatment at 2050~F restored both ultimate and yield strengths (Table 3, Spec. R2-14 and Spec. R2-15). Ductilities, however, may not have been completely restored. Because no creep was involved, any inability to completely restore ductility must have resulted from some difference in structure before and after re-heat treatment and not to a permanent damage arising from creep. Complete Re-Solution after Exposure to Creep at 1800~F The ability of the "A" treatment to restore thermally-induced changes after 100 hours unstressed exposure at 1800~F was demonstrated on specimens R.-ll and R-201. Creep-exposure for the same time period also had a severe effect on the tensile properties (Table 3, Spec. R-26 and Spec. R-45). An indication of the effect of the "A" treatment in restoring properties after creep-exposure was obtained by exposing specimens to 3.46 and 28. 4 percent of creep in 100 hours at 1800~F prior to re-heat treatment. These specimens were then heated at 2150~F for 2 hours, air cooled, reheated to 1975~F for 1 hour, water quenched, and aged at 1400~F for 16 hours. Both specimens had about the same ultimate and yield strengths as the specimen (R-141) simply re-heat treated. The ductility of-Specimen R.-124 with 3. 46-percent creep may have been slightly reduced and was definitely reduced for Specimen R-116 with 28. 4-percent creep. The restoration of properties after 28. 4-percent creep was, however, remarkably effective and indicated that if there was permanent damage from creep at 1800~F it was very small. Influence of Temperature and Time of Creep Exposure on Response to Heat Treatment The experiments involving material exposed at 1800~F for 100 hours demonstrated that a complete re-solution essentially restored tensile properties at room temperature when the exposure conditions were the most severe used in this investigation. Presumably, any less drastic exposures would also respond to a similar heat treatment. The data did suggest, however, that information could be obtained on the structural changes reducing properties from less drastic exposure if the re-heat treatment was limited to gamma prime and M23C6 re-solution by using the original "R" heat treatment of 1950~F solution and 1400~F aging. A group of specimens were exposed for 10 hours at 1400~, 1600~, 1700~, and 1800~F under stresses producing from 3 to 6 percent of creep (Table 4). This exposure generally was about 70 to 80 percent of the 17

estimated rupture life. The specimens were then re-heated to 1950~F for one-half hour, air cooled, aged at 1400~F for 16 hours, and finally re-machined. The creep curves for each exposure are plotted in Figures 9-11 together with bar graphs indicating the level of room temperature tensile properties at each stage of the exposure and re-heat treatment cycle. Summary plots to illustrate the effects of exposure temperature and time are included in Figure 12. Properties were essentially unchanged from the original condition (Table 4) for specimens exposed at 1400~ and 1600~F (Figs. 9, 10, and 12). The increase in yield strength from the Bauschinger effect was removed by re-heat treatment. The loss in ductility at 1600~F was restored. The ultimate and yield strength reductions from exposure to 1700~ and 1800~F were restored, but the accompanying losses in ductility were not (Figs. 10, 11, and 12). When the exposure time was increased to 100 hours, the re-heat treatment only returned both strength and ductility to near original values for exposure at 1400~F (Figs. 9 and 12). Strengths were restored after exposure at 1600~F but not ductility (Figs. 10, 12). Yield strength was restored after exposure at 1800~F but not ultimate strength or ductility (Figs. 11, 12). As previously discussed, the re-heat treatment apparently had to be 2150~F to include re-solution of M6C carbides for exposure for 100 hours at 1800~F (Fig. 12). Presumably, this temperature would also be required to restore ductility after 100 hours exposure at 1600~F. The results for either time of exposure (Fig. 12) showed no appreciable influence of creep on the ability for property restoration by re-heat treatment. Summary of Tensile Property Studies The results indicate that deterioration of tensile properties during prior creep exposure involves the following processes: (a) Any changes in yield strength can be substantially restored by a treatment at 1950~F. When exposures are at a sufficiently low temperature to introduce Bauschinger effects or strain hardening, the reheat treatment provides stress relief. When the temperature and time of exposure are high enough to cause reductions in yield strength due to loss of the original fine dispersion by agglomeration of gamma prime, re-solution of the gamma prime at 1950~F and re-aging restores yield strength. 18

(b) The initial losses in ductility for short time or low temperature of exposure apparently can be restored by a treatment at 1950 F. Either these are due to removal of M23C6 carbides which are redissolved at 1950~F; the treatment at 1950~F changes carbide distribution so that it does not effect ductility; or strain hardening is relieved. (c) When the exposure temperature is increased above 1600~F and time of exposure is increased beyond 10 hours, ductility cannot be restored by re-heat treatment at 1950~F (Fig. 12). Because treatment at 2050 "F was not effective in one case while 2150"F was effective in all cases (Table 4), it appears that the re-solution treatment must be high enough to dissolve M6C carbides when they have formed or reached a given size or dispersion. (d) Ultimate strength will not be restored by gamma prime resolution when this treatment does not restore the more severe cases of reduced ductility. (e) Creep in itself did not appreciably change the response to reheat treatment. The only effect it might have would be to accelerate the structural changes previously discussed. (This is based on specimens with 3 to 6 percent of creep in 10 and 100 hours at 1400~ to 1800~F. One case of 28-percent creep at 1800~F also responded. ) The indications are that internal cracking was not occurring during creep; or, if it did occur, it was not sufficient to noticeably effect properties. In considering this general summary, it is important to again recognize that surface damage from prior exposure and during re-heat treatment was eliminated byre-machining the specimen surfaces. If this had not been done, the degree to which ultimate strength and ductility were restored would have been considerably less. RE-HEAT TREATMENT AND CREEP-RUPTURE PROPERTIES Differences in the plastic flow mechanism between room temperature and the creep-exposure temperatures could be responsible for the ability to restore the room temperature tensile properties by re-heat treatment after creep. That is, permanent damage from prior creep might have little effect on tensile properties due to differences in the mechanism of plastic flow. Accordingly, it appeared that a study of the effect of re-heat treatment on creep-rupture properties after creepexposure for part of the rupture life would provide information further delineating the principles governing creep damage. This procedure also would extend the measure of creep damage by assessing it in terms of 19

properties quite different than the mechanical properties at room temperature. The information could have considerable practical utility in those cases where such re-heat treatments would be feasible and surface damage effects could be controlled. It is generally considered that creep-rupture life which has been used up cannot be restored by re-heat treatment. When such treatments are successful in prolonging creep-rupture life, structural deterioration reducing creep-rupture life presumably has been restored. The reason for the permanency of creep-damage to subsequent creep-rupture life has never been fully explained. The explanation based on microcracking has appeared to be deficient. For interrupted creep-rupture tests, the rupture time after reheat treatment should be no longer than the difference between the expected uninterrupted (i. e., "normal") rupture life and the time at which the test was interrupted. (This is an expression of the so-called "life fraction rule". ) Interrupting and re-starting a test on a re-heat treated andre-machined specimen involves an adjustment of stress which might increase the rupture time slightly. Also, the re-machining could increase life if surface attack was an important factor controlling rupture. However, any marked increase in rupture life would presumably be due to the restoration of a structural change which was markedly reducing creep resistance with time of testing. Re-Heat Treatment by Re-Solution of Gamma Prime and M23C6 Carbides Only A group of specimens were exposed to creep at 1400~, 1600~, 1700~, and 1800~F generally for 10 or 100 hours. The stresses were such that the time periods represented 60 to 90 percent of the expected rupture life (Table 5). The initial "R" heat treatment of one-half hour at 1950~F, air cool, plus aging at 1400~F for 16 hours was then reapplied. After remachining to eliminate surface damage and confine the results to creep damage to the internal structure, the specimens were returned to creeprupture testing under the original conditions. Because the loads were changed to give the same stress as the initial stress, any stress intensification from reduction of area during the initial exposure was eliminated. In every case (Table 5 and Fig. 13), the rupture life after re-heat treatment was considerably longer than the additionaltime expected for rupture to occur in uninterrupted testing. It ranged from 1. 82 to 4. 60 times as long as expected on the basis of no restoration of creep life ("life fraction" rule) (Fig. 14a). This measure is, perhaps, not as im20

portant as the fraction of expended life recovered by the heat treatment. The percentage recovered tended to decrease with increasing test temperature and time (Fig. 14b). The specimen exposed for 10 hours at 1600~F apparently completely recovered its creep-rupture life during the re-heat treatment. The specimen exposed for 10 hours at 1400~F, while it recovered a substantial amount of life, did not reach 100 percent. It did, however, rupture with low ductility (Fig. 14d) and it is suspected, for this reason, that something may have been wrong with the test. (This specimen also represented the maximum expenditure of life during the first exposure of any of the specimens studied. ) The general trend in Figure 14b for decreasing recovery with increasing temperature would be in agreement with the general trend of the data for the effects on properties at room temperature (Fig. 12). On the other hand, if the high recovery tests are due to erratic test behavior, then all the specimens were damaged to the extent that only about 50 percent of the original life was restored by the re-heat treatment (Fig. 14b and 14c). The elongation during the first creepLexposure plus that during the second generally totalled more than for uninterrupted testing. They were as high during the second exposure as for uninterrupted testing, except for prior creep at 1400~F (Fig. 14d). The re-heat treatment initially returned the creep resistance very nearly to that of the material at the start of the first exposure. Examination of the creep curves (Figs. 15 and 16), however, indicates that the creep rates increased with time during the second exposure faster than during the first. This tendency was more apparent the higher the temperature of the first exposure (Fig. 14e). Some factor inherited from the first creep exposure affected creep after re-heat treatment. From these data, it appears that: (1) Complete recovery of creep-rupture life by re-heat treatment andre-machining after exposure at 1400~F was prevented by incomplete recovery of ductility. (2) When the exposure temperature was 1600~F, nearly 100-percent recovery was obtained because both creep-resistance and ductility were recovered. (3) The inheritance factor preventing complete recovery of prior creep-exposure at higher temperatures appeared to be the faster rate of increase of creep with testing time. 21

Complete Re-Solution and Re-Heat Treatment The response to re-heat treatment at 1950~F was somewhat similar to that observed in tensile tests at room temperature. This suggested that solution was incomplete and that a full solution treatment would completely restore creep-rupture life. Accordingly, specimens exposed to creep at 1800~F were heated at 2150~F for 2 hours and air cooled. Like the specimens used for tensile tests, a treatment of 1 hour at 1975~F, water quench and age at 1400~F for 16 hours was then applied. (This sequence was designated the "A" treatment) After re-machining, the specimens were returned to rupture testing at 1800~F under the same stress as the first creep exposure. The rupture times were longer than those expected for uninterrupted testing of the original material (Table 5 and Fig. 17). In other words, the complete solution treatment at 2150~F not only resulted in complete recovery from the prior creep but also added more rupture life. In view of this result, the influence of the treatment on specimens which had not been exposed to creep was determined. The rupture times were the same as those for the specimens which had been exposed to creep (Fig. 17). From the viewpoint of rupture life, there was no residual effect from creep after complete re-solution. The ductility during the second creep-exposure was about the same as for specimens heat treated at 2150~F without prior creep (Table 5). The creep curves were virtual reproductions of those for similarly heattreated specimens which had not been exposed to creep (Figs. 18 and 19). Apparently, the inclusion of a complete re-solution at 2150~F in the re-heat treatment resulted in complete recovery from prior creep. The inheritance factor after re-heat treatment at 1950~F of a more rapid increase in creep with time of testing was eliminated. Because the testing was limited to prior creep at 1800~F, it is not known if the reduced ductility found after creep at 1400~F and re-heat treatment at 1950~F would also have been eliminated. Influence of Surface Damage on Response to Re-Heat Treatment Part of the prolongation of rupture life from re-heat treatment could be due to the removal of surface damage by re-machining, In order to obtain some information on this point, two tests were carried out. Limitation of time and funds prevented a more thorough study. 22

Specimen R-183 with the original heat treatment at 1950~F was exposed to creep at 1600~F for 10 hours to 2. 5-percent creep strain. After re-heat treatment at 1950~F and aging at 1400~F, it was returned to creep-rupture testing without re-machining. The rupture time was nearly the same as for the specimen which was both re-heat treated and re-mnachined (Table 5). Ductility was lower and was probably responsible for a slightly shorter rupture life. Creep resistance appeared to be similar to that during first exposure (Fig. 20). Specimen R-187 was similarly exposed to creep, re-machined but not re-heat treated. The rupture time during the second exposure was 60 percent of that for the specimen both re-heat treated and re-machined. This did, however, represent a considerable prolongation of life beyond that of an uninterrupted test. The prolongation of life was apparently due to the increased elongation after re-machining. Creep resistance was somewhat reduced (Fig. 20). These very limited data suggest that re-heat treatment restores creep resistance while re-machining is required to restore ductility. The indication is based on only two tests limited to prior creep at 1600~F. It was previously shown that re-machining presumably did not restore ductility when the exposure temperature was 1400~F and the re-heat treatment was limited to 1950~F. If these results are indicative of the magnitude of the two effects, re-heat treatment is considerably more effective in prolonging rupture life than re-machining. This is possibly related to the shape of the creep curves. The restoration of ductility by only re-machining operated when the creep resistance had already deteriorated to a considerable extent. The rapid creep used up the ductility in less time. General Principles from Study of Re-Heat Treatment on Creep-Rupture Properties In general, re-heat treatment and re-machininghad the same effects on future creep-rupture properties as it did on tensile properties at room temperature. A re-solution treatment at 1950~F followed by aging at 1400~F to restore the gamma prime dispersion, did not completely restore properties. Complete re-solution by including a re-heat treatment at 2150~F did completely restore properties. In fact, this treatment actually increased the creep-rupture life beyond that expected for the original material heat treated at 1950 F. The inheritance factors when the re-heat treatment was limited to 1950~F were apparently somewhat different than for tensile properties. 23

When the first creep-exposure was at a relatively low temperature (1400~F), ductility during the re-exposure to creep was reduced. After higher temperatures of exposure to creep, ductility was restored but the creep rates increased with time faster than during the original exposure. The tensile properties at room temperature could be restored after creep exposure at the lower temperatures but not after the higher temperatures. The latter was due to non-recovery of ductility. Possibly, the same structural feature, incomplete solution of M6C carbides, was responsible for both effects. The cause for the inability to recover creep-rupture ductility after creep exposure at the lower temperatures is not clear. Possibly, the M6C formed affects creep-rupture ductility but not tensile at room temperature. The data suggest that restoration of gamma prime by re-heat treatment restores both yield strength at room temperature and creep-resistance. The removal of surface damage by re-machining restores ductility in creep-rupture tests. Due to concurrent increases in creep resistance from re-heat treatment, this was less effective in prolonging creeprupture life than it was in restoring properties at room temperatures. In other words, the restoration of creep-resistance apparently was the predominant factor. The absence of permanent loss of creep-rupture life by prior creep indicates that the principal damage factors during prior creep were structural changes. The potential creep resistance of the material is such that considerable creep extension introduces little damage if the structure is restored. Certainly, there was no result suggesting that internal cracking was involved. The difference in plastic flow between room temperature tensile tests and creep-rupture conditions only reflected itself in the differing ways the structures control the two properties when the re-solution during re-heat treatment was incomplete. Table 5 includes tabulations to show the degree by which re-heat treatment prolonged rupture life and changed total ductility. These values were obtained by summing the respective values for the exposure to creep before and after re-heat treatment. It should be recognized that these values cannot be considered for practical applications without consideration of the effects of surface damage and the possibility of re-machining, The latter is frequently impossible. Surface damage can have serious deleterious effects on other properties than those considered. In addition, such factors as warpage during re-heat treatment could be serious. The investigation also did not consider possible deleterious effects of a 2150~F treatment on other properties, particularly at temperatures below 1800~F. 24

MICROSTRUCTURAL ASPECTS OF CREEP DAMAGE As has previously been indicated, the experimental program was based, in many ways, on the expected effects of creep-exposure on the microstructure of Rene' 41 alloy. In particular, the following features were used: 1) The main strengthening feature of the microstructure was the precipitation of Ni3(Al, Ti), commonly referred to as gamma prime. Alloys containing gamma prime are known to undergo losses in strength as it agglomerates during exposure at elevated temperatures. The previous report (Ref. 1) showed electron micrographs which indicated qualitatively that yield strength decreased as the size of the gamma prime particles increased with exposure temperature and time. This feature of damage from exposure has now been examined in more detail by making quantitative measurements of the gamma prime as a function of exposure conditions. 2) The role of gamma prime was studied further by determining the influence of re-heat treatment to dissolve agglomerated gamma prime and reprecipitate it in the original dispersion. This was done for both mechanical properties at room temperature and subsequent creeprupture properties. Studies of the influence of re-heat treatment on microstructure were then carried out. 3) One of the objectives of the re-heat treatment program was to determine if there was permanent damage from creep which could not be restored by heat treatment. In particular, the possibility of microcracking as a source of permanent damage was examined microscopically. 4) Other research on nickel-base gamma prime-strengthened alloys had demonstrated that carbides precipitated at grain boundaries could be damaging to ductility (Refs. 11, 12). In this research, it was previously shown (Ref. 1) that reduced ductility at room temperature was associated with the growth of fairly massive carbides in the grain boundaries during exposure. X-ray diffraction analyses of extraction residues from exposed samples had shown that both M23C6 and M6C carbides could be present in material with either high or low ductility. The choice of re-heat treatment temperatures for the present studies was based on the restoration of gamma prime by heating at 1950~F but not re-solution of M6C carbides. The choice of 2150~F was based on the expected re-solution of M6C carbides. The actual structures and types of carbides after exposure and after re-heat treatment were accordingly investigated for this report. 25

Influence of Exposure and Re-Heat Treatment on Microstructure The series of photomicrographs showing the coarsening of gamma prime with increasing exposure temperature for 10 and 100 hours from Reference 1 have not been repeated for this report. The original optical microstructure (Fig. 21) shows some grain boundary precipitation and titanium carbonitrides. The fine gamma prime precipitate present is not evident optically but is shown by the electron micrograph of Figure 50. The exposures up to 1400~F did not noticeably change the microstructure. As shown in the previous report, starting at 1500~F, exposure at higher temperatures coarsened the gamma prime and increased the amount of precipitate in the grain boundaries. Figures 22,40, 43 and 44 show the extreme changes developed by 100 hours at 1800~F Figures 65 and 67 show the influence of prolonged exposure at 1700~F. The gamma prime particles became large enough so that they were evident in the optical microstructure. The grain boundary precipitates became continuous. As shown in Reference 1, both massive gamma prime and carbides formed in the grain boundaries. Exposure at 1600~ to 18000F would be expected to result in solution of gamma prime as required by the phase relationships. This took place and as a result, during cooling, fine gamma prime re-precipitated in the matrix between the larger undissolved gamma prime particles as is shown by Figures 40, 43, 44, 65, and 67. In the previous report, it was shown that this fine re-precipitation did not always occur under conditions and for reasons which have not been established. The re-heat treatments at 1950 F dissolved the gamma prime and on re-aging developed about the same gamma prime dispersion as the original material. Figures 25 through 33 show, however, that as the carbides in the grain boundaries increased with creep-exposure temperature and time, they remained in the grain boundaries in increasing amounts after re-heat treatment. Apparently, this is the reason why re-heat treatment, simply involving re-solution of gamma prime, was not always successful in restoring ductility at room temperature. (The effects on mechanical properties are shown in Figure 12. ) These micrographs show that massive gamma prime was dissolved from the grain boundaries and, therefore, was not involved in the loss in ductility. (It should be noted that the electron micrographs of these figures do not show gamma prime reprecipitated. This was due to the low magnification necessary to properly show the carbides. ) Re-heat treatments involving complete re-solution at 2150~F were, however, effective in reducing the grain boundary carbides (Figs. 23, 24, 34, 35). 26

Relation of the State of Gamma Prime after Exposure to Yield Strength The qualitative observations between gamma prime particle size and yield strength after exposure (Ref. 1) suggested the possibility that quantitative measures of gamma prime would provide a useful means of defining the general principles governing creep damage. Accordingly, measurements were made to define the volume fraction of gamma prime, its particle size, and spacing as a function of exposure conditions with the results given in Table 6. The method used was to carry out lineal analyses on electron micrographs, using the adaptation developed by Rowe (Ref. 13) to correct for the particle size being smaller than the etching depth. These structural parameters were then correlated with the yield strength. In making these measurements, only samples exposed without stress were used to avoid the complication due to residual stresses influencing yield strength. The samples given prolonged exposure in the dilatometer were added to the study. In attempting to measure specimens exposed up to and including 1400~F, it was found that the change in gamma prime for the exposure conditions considered was so slight that the time-consuming measurements would not show a significant change. This is quite certainly the reason that there was no change in yield strength after these exposures. The measurements of structures exposed at 1600~, 1700~, and 1800~F gave the following results: 1) The average particle size increased with both temperature and time of exposure (Fig. 36). 2) The volume fraction of gamma prime decreased with exposure time with a strong tendency towards a constant minimum value (Fig. 37). The objective of these measurements was to relate the state of the gamma prime to the yield strength at room temperature after exposure. Both the increase in size of gamma prime particles and the decrease in volume would commonly be expected to reduce yield strength. The scatter in data and lack of time to conduct check and additional tests where desirable made it necessary to analyze the results in terms of general trends. Within this limitation, it appears from the data that: 1) The yield strengths at room temperature (Fig. 3) tend to approach a common level with exposure time. The common level is attained in shorter times with increasing temperature of exposure. 2) The particle size of gamma prime increases with both exposure temperature and time. 27

3) The volume fraction of gamma prime tends to approach a common level with exposure time. This is attained in somewhat less time with increasing exposure temperature. In other words, the yield strengths tend to level out along with the volume fraction of gamma prime while the particle size continued to grow with exposure time. The time required for this condition decreased slightly with increasing exposure temperature. Particle size, however, continued to increase with both exposure temperature and time. An extensive study of the explanations for strengthening from dispersed phases was carried out with the aid of the literature survey of Bunshah and Goetzel (Ref. 14). The treatment developed by Meiklejohn and Skoda (Ref. 15) from studies of iron particles dispersed in solid mercury seemed to best explain the results. Moreover, their treatment seems best substantiated from a theoretical viewpoint (Ref. 16). The experiments of Meiklejohn and Skoda showed that strengthening from a dispersed phase -- if the phase is truly a dispersion -- depends on the ratio of h/\h, where h is the particle diameter and A is the interparticle spacing. In turn, from geometric considerations, it can be shown that h/A is a function of the volume fraction of the dispersed phase; namely, h/A is proportional to f. Thus, the yield strength is a function of the volume fraction of the precipitate, i. e., ry-o =. 8-f-, where -y is the strength with the dispersion, o-o is the strength of the matrix, and "a" is a constant. However, if the yield strength, at a constant volume fraction, I, does vary with the particle size, h, then the foregoing conditions are not being met. Meiklejohn and Skoda explain this on the basis of coherency effects, and show that in the case of coherency, an effective volume fraction, f', must be used; defined as the stressed region through which dislocations will not pass. Thus, the operative effect is the blocking of dislocations by the actual or effective dispersion. In applying these considerations to the present data, a plot of -ry versus h (Fig. 38) showed some dependence of yield strength on particle size -- although the data did not permit comparison to be made at constant volume fraction. Thus, the conditions for a rtrue" dispersion were not entirely met. With this in mind, the yield strength data were all tested with the Meiklejohn-Skoda volume fraction parameter, O. 82-ft, with the results being plotted in Figure 39. Allowing for experimental scatter, the correlation appears to be good, and in particular, shows excellent agreement with one facet of the Meiklejohn-Skoda equation in that the intercept, 28

-o, at zero volume fraction, is the experimentally determined yield strength of the matrix (mill annealed) material prior to aging. The correlation of Figure 39 thus suggests that the volume fraction, f, was the strong variable governing yield strength, with the particle size effect suggested by Figure 38 of secondary importance. The break in the curve, or limiting yield strength, indicated in Figure 39 for large volume fractions was also reported by Meiklejohn and Skoda for their iron-mercury dispersions. Their explanation for the effect was that at the maximum stress either cross-slip occurred with dislocation loops climbing over the particles or that dislocations passed through the particles. For precipitation hardening alloys, Meiklejohn and Skoda suggested that after attainment of maximum strength, as the particles continue to grow, a decrease in yield strength is expected until the particles attained a certain size. At this point, the strength would be expected to level out with increasing particle size and become only a function of volume fraction. The general trends of the data of Figure 38 show a rather remarkable similarity to this prediction and suggest that: 1) The yield strengths initially decreased when the particles of gamma prime were small due both to the increase in particle size and to the decrease in volume fraction. 2) The yield strengths levelled out and became virtually independent of exposure time and temperature because the particles of gamma prime attained a size where volume fraction was the controlling feature and the volume fraction attained nearly a constant value independent of the exposure temperature. 3) Yield strengths did not continue to decrease with exposure time (Fig. 3) because the gamma prime particles attained a size beyond which size changes had no further effect. 4) The variations in yield strength with exposure temperature and time for exposures of 10 to 100 hours reflect the combined effects of rate of increase in size of gamma prime particles and the rate of decrease in volume fraction as they were influenced by exposure conditions before the limiting conditions were attained. Extensive attempts were made to utilize the data to correlate the properties in other terms of the fundamental characteristics of the particle size and interparticle spacing. The literature is not clear on the proper way to do this and attempts to correlate the data in this manner 29

led to a number of fair correlations and others which did not correlate. It is believed, however, that more time to obtain more data in the proper ranges of properties and particle size together with proper recognition of the limiting size effect would have resulted in considerable progress in further clarifying the factors controlling yield strength of gamma prime-strengthened alloys. The as-treated yield strength of material aged at 1400~F ("R" condition) was higher than that for material aged at 1650~F (R2"11 condition). The data of Figures 38 and 39 indicate that this was due to both the larger volume fraction and smaller particle size of the gamma prime. Both heat treatment conditions rapidly attained the same yield strength as a result of exposure because both of the structural parameters rapidly attained the same value. There are a number of features of the data which limit the generality and reliability of the explanation proposed: 1) The decrease in measured volume fraction of gamma prime with exposure is not understood. Rowe (Ref. 13) encountered the same effect in M252 and Inconel 700 alloys. Bigelow (Ref. 17) had also reported this phenomena. Bigelow proposed that it was due to the entrapment of matrix atoms in the gamma prime which during exposure had an opportunity to diffuse out of the gamma prime. Rowe was unable to completely explain his results on this basis and suggested the possibility that coherency effects where particles were small may have somehow resulted in the measured volume fraction being increased beyond the true volume fraction. In all cases, the initial measured volumes were larger than would be predicted from known phase relations. 2) The measurements of gamma prime were confined to the particles freely dispersed in the matrix. The massive gamma prime accumulated in the grain boundaries was not counted. As Rowe (Ref. 13) observed, the gamma prime in the grain boundaries was, however, nowhere near sufficient to account for the decrease in measured volume. 3) In this investigation, the exposure conditions and phase relationships required solution of gamma prime from that present after aging in accordance with the increasing solubility with temperature. In some cases, this resulted in reprecipitation on slow cooling in the form of fine particles in the matrix between the larger particles of undissolved gamma prime. In the measurements, the unsatisfactory procedure of merely averaging the particle size was used. In other samples, the fine precipitate was absent. Time did not permit clarification of these two effects. It must be recognized, however, that part of the measured 30

volume decrease could be due to re-solution of gamma prime. The attainment of a volume fraction nearly independent of exposure temperature suggests that this was not much of a factor. 4) Meiklejohn and Skoda (Ref. 15) treated the case where the dispersed phase was hard in a soft matrix and was chemically and structurally widely different from the matrix (i. e., little opportunity for coherency). It must be recognized that in alloys such as Rene' 41 gamma prime is very similar to the matrix (gamma phase) in both crystal structure and hardness. Thus, the present conditions are widely different than were postulated by Meiklejohn and Skoda. 5) The particle size at which the strengthening from the dispersed phase became a function only of volume fraction in Meiklejohn and Skodas' work (Ref. 15) was smaller than the point of apparent levellingoff in this investigation (Fig. 38). The latter data are not, however, too definite due to the difficulty of accounting for the volume fraction and particle size varying simultaneously. In addition, data scatter could have played a part. There seemed to be a tendency for higher strength for smaller particle size in the range where volume fraction was decreasing. It is possible that this apparent discrepancy was due to the similarity of gamma prime and gamma and to the retention of coherency to larger particle size for gamma prime. Effect of Exposure on Matrix Lattice Parameter Measurements of the lattice parameter of the matrix showed a decrease (Table 7) with increasing temperature and time of exposure. Calculations of shrinkage based on these measurements agreed with those measured in the dilatometer. Table 7 also includes measurements made by Beattie (Ref. 11) which show good agreement between the two sets of measurements when allowance is made for the lower solution temperature and aging at 1400~F for the present investigation. These data suggest that the volume shrinkage was due to precipitation removing odd-sized atoms from solution. In this case, this would presumably mainly be carbon and molybdenum. There are, however, uncertainties involving the role of the apparent volume fraction change of gamma prime. If the apparent volume change of gamma prime was due to some preferential segregation of other atoms not involving size differences, there could have been a shift in matrix lattice parameter and volume. It should be noted, however, that shrinkage continued after the volume fraction of gamma prime apparently became nearly constant. The indications are, therefore, that precipitation of odd-sized atoms as carbides was the major reason, even though the amount of shrinkage seems large for this to be the sole cause. 31

Relationship of Carbides to Properties The microstructural studies of carbides were qualitative in nature. Attempts were made to make this more definite. As was shown in Reference 1, x-ray diffraction analysis of extraction residues after exposure showed both M23C6 and M6C carbides present when ductility was high as well as when it was low. Additional diffraction data for the dilatometer specimens given prolonged exposure at 1600~ to 1800~F are included in Table 13. The main correlation of properties with carbides seemed to be the deterioration in ductility as the total amount of carbide increased. Other investigations (Refs. 11, 18) had shown that M23C6 tends to break down to M C as the exposure temperature and time are increased. Moreover, it ha: been found that M23C6 redissolved between 1700~ and 1800~F. Because M23C6 had been found in samples exposed as high as 1900~F for 10 hours, this was difficult to understand. Further information on this point was obtained by carrying out x-ray diffraction analysis of extracted carbides after re-heat treatment (Table 8). After exposure for 100 hours at 1800~F, both M23C6 and M6C were present in Specimen R-104 (see also Fig. 22). (The TiC was originally present and remained constant with re-heat treatment and apparently was inactive.) After re-heat treatment at 2150~F for 1 and 2 hours (Figs. 23 and 24), the patterns for M23C6 and M6C became weaker. The two-hour treatment reduced the M23C6 and M6C to quite small amounts with possibly more M23C6. These results again confirmed that treatment at 2150~F was effective in dissolving carbides (compare Figs. 33 and 35). In this case, the carbides found could have formed during aging at 1400~F. This also suggests that they may have been present in exposed samples as a result of opportunity to form during cooling. Considerable effort was expended to produce extraction replicas to isolate the carbides in place in the microstructure and then use electron diffraction for identification purposes. The specimens used had been exposed at 1800~F for 100 hours (Fig. 40). When the technique of leaving the gamma prime in the structure before extraction of the carbides was used, the replicas illustrated by Figures 41 and 42 were obtained. The dark areas are the extracted carbides. The very light areas adjacent to the carbide in Figure 41 were due to a tear in the replica, a major difficulty with the technique. In Figure 42, the massive gamma prime adjacent to the carbide can still be seen. When the gamma prime was etched out prior to applying the extraction replication procedure, the replicas shown by Figures 45 through 48 were obtained. These included areas of what appeared to be thin films of carbide. Normal electron micro32

graphs of collodion replicas of this sample are shown in Figures 43 and 44. Satisfactory electron diffraction patterns of these, characteristic of thin films (Fig. 49), were obtained. The patterns indexed as M23C6. Useful patterns could not be obtained from the more massive carbide particles which were extracted since they were not suitable for electron diffraction purposes. It should be recognized that the thin films of M23C6 could have formed during cooling after exposure and may not be entirely characteristic of prolonged exposures. To summarize the carbide studies, the decrease in M23C6 and M6C carbides during re-solution treatment at 2150~F was substantiated. Extraction replicas further identified carbides in exposed samples. The association of reduced ductility with massive carbides was, therefore, given further substantiation. The simultaneous presence of M23C6 after exposure and after heat treatment was proven by additional data. The occurrence of M23C6 as thin films was shown. The reason why M23C6 was found after exposure or heat treatment above the reported solution temperature of 1700~ to 1800~F was not determined, although it probably was formed during cooling after exposure or during aging after re-heat treatment. The presence of M23C6 as thin films would be consistent with its formation under those conditions. Increasing the re-heat treatment temperature to 2150~F did result in extensive carbide re-solution (Figs 23, 24, 34, and 38). The carbides after 2 hours at 2150~F were even less than in the original condition of heat treatment (Figs 50 and 51). The re-solution of carbides, therefore, did coincide with the restoration of ductility. The reprecipitated gamma prime particles (Fig. 51) were finer than when the material was heat treated at 1950~F (Fig. 50) This coincided with the increase in creep-rupture strength when the re-heat treatment temperature was raised to 2150~F from 1950~F, Additional Observations from Structural Studies A relationship between structure and the inheritance factor of more rapid increase in creep rate with time after re-heat treatment at 1950~F than in the first exposure was not identified. As discussed previously (page27), there were volume differences in measured gamma prime as well as changes in particle size during aging and exposure. It might be that if such measurements had been made, it would be found that the volume of gamma prime was less whenever carbides were not redissolved and that re-solution of carbides increased the volume percentage of gamma prime. Presumably, the re-solution of carbides at 2150~F was responsible for the reduced size of the gamma prime particles after aging, although 33

other factors associated with nucleation and growth could have been involved. One of the most important results of the microstructural examinations was the absence of microcracking after creep in any of the samples examined. The authors have found that microcracking is a prominent feature of creep in all other nickel-base gamma prime-strengthened alloys with which they have had experience. At present, this absence of microcracking appears to be unique to Rene' 41 and may be the reason why creep damage was negligible. EFFECT ON TENSILE PROPERTIES OF SURFACE REACTIONS DURING EXPOSURE Although the previous portion of this investigation (Ref. 1) was concentrated on r-machinedspecimens, a few tensile tests were conducted on specimens which were not machined to remove surface damage after exposure. The results suggested that the reduction in ultimate strength and ductility at room temperature from surface damage was similar in magnitude to that due to structural changes. Additional data have now been obtained for exposure without and with creep. Exposure Without Creep The exposure conditions were 10 hours at 1200~ to 1800~F and 100 hours at 1000~ to 1800~F without stress The tensile properties (Table 9) and the graphical comparisons (Figs. 52 and 53a) to specimens with re-machined surfaces indicate the following: 1) Yield strengths (Figs. 52 and 53a) were the same for both remachined and as-exposed surfaces. 2) After exposure for 10 or 100 hours, the as-exposed surfaces resulted in lower ultimate tensile strength and ductility (Figs. 52 and 53a) than for the re-machined specimens for exposure at 1400 F and higher temperatures. The degree of reduction increased somewhat with exposure temperature. An exception was exposure for 100 hours at 1800 F where the limited data indicate little difference between as-exposed andre-machined specimens. Exposure at 1200~ and 1300~F did not result in reduced properties from surface effects. 3) The amount the as-exposed surfaces reduced properties below those of the machined specimens (Figs. 52 and 53a) was essentially the same for either 10 or 100 hour s of exposure at a given temperature, 34

except for 1800~F. 4) Specimens with as-exposed surfaces exhibited reduced ductility after exposure for 10 hours above 1300~F, while the re-machined specimens exhibited no change (Figs. 52 and 53a). The tensile strength was also reduced below that due to structural changes. When exposed for 100 hours, the surface and structural change effects were approximately equal. The exception was 1800~F when the structural changes appeared to be responsible for nearly all the reduction. Exposure With Creep Specimens were exposed to creep for 10 hours from 1200~ to 1800~F and for 100 hours at 1200~ and 1300~F and then tensile tested at room temperature with as-exposed surfaces. The data obtained (Table 9) indicate: 1) The introduction of creep caused the as-exposed surfaces to reduce ductility for exposure at 1200~ and 1300~F (Fig. 54). This did not occur for specimens exposed without creep (Fig. 52). Specimens re-machined after creep exposure also showed no loss in ductility, except when deep cracking occurred. As will be shown in the next section, this was due to an extraneous effect. 2) The surfaces as exposed to creep at 1200~ and 1300~F had no effect on yield or ultimate tensile strengths (Fig. 54). The reduced ductility did not affect ultimate strength since the residual effects of creep which manifested themselves as the Bauschinger effect and strainhardening caused a change in the shape of the stress-strain curve such that the yield strength and ultimate strength were almost equal. 3) The ductility after exposure above 1300~F for 10 hours was essentially constant at each temperature for specimens with as-exposed surfaces (Figs. 55 and 53b) with increasing amounts of creep. The damage from exposure without creep was the same as when creep occurred. Ultimate tensile strengths were also reduced by as-exposed surfaces. Creep had little effect except possibly to increase damage at 1800~F. At 1400~ and 1600~F, it may have reduced the surface damage. General Effects from As-Exposed Surfaces Surface alterations which can include such things as surface cracking, oxidation, decarburization, and de-molybdenumization reduced ultimate tensile strength and ductility at room temperature after exposure. 35

Surprisingly, the only significant effect from creep was reduced ductility for the 1200~ and 1300~F creep-exposures. The damage was not significantly increased by increasing exposure times from 10 to 100 hours. There was some increase from increasing temperature of exposure. The mechanism of damage, therefore, appears to be fairly complicated because surface reactions certainly increase both with time and temperature of exposure, and creep would be expected to intensify cracking. Deep Cracking During Creep at 1200~ and 1300~F Exposure to creep at 1200~ and 1300~F resulted in progressive reduction in ductility with increasing creep up to about 4-percent strain (Fig. 54). For larger amounts of creep, there was a marked increase in the loss in ductility along with a tendency for reduced ultimate strength where the ductility was particularly limited. These specimens exhibited oxidized fracture surfaces (Fig. 56) extending in from the surface to varying degrees, indicating that cracks had penetrated during exposure deeper than the 0. 025-inch machined off before testing. This phenomena has been designated "deep cracking". Measurements of the depth of deep cracking are summarized in Table 11 and plotted as a function of prior strain in Figure 54. Re-examination of these specimens disclosed that the deep cracks were associated with points of thermocouple attachment. When this was noted, three specimens were exposed at 1300~F with the thermocouples attached only to shoulders. Even when the creep strain was 5.4-percent there was no evidence of deep cracking (Table 11 and Fig. 54). The ductility values were in accordance with a prolongation of the curves for smaller creep strains where there had been no deep cracking. Further examination showed that the cracks generally originated at a point about one-third around the circumference from the point where the thermocouple beads were in contact with the specimen. This is the point at which the wire fastening the thermocouple to the specimen first comes in contact with the specimen. The cracks as revealed by dye-penetrant testing are shown by Figure 56. Further testing (Tables 9, 11 and Fig. 54) showed that the condition developed most extensively when alumel wire was used to attach the thermocouple. When chromel was used, the penetration was reduced. Specimen R-152 was exposed to 5. 6-percent creep at 13000F with a loop of alumel wire wrapped around one end. When tested at room temperature, deep cracking which had occurred where the alumel loop was attached, had reduced ductility (Fig. 54). One-half of the gage length of Specimen 36

R-156 was wrapped with alumel and the other half with chromel before exposure to 5. 9-percent creep in 10 hours at 1300~F. Dye-penetrant examination showed numerous cracks on the end wrapped with alumel, with practically none on the chromel half. When 0. 046-inches was machined off the surface and the specimen then subjected to tensile testing, the properties (Fig. 54) were in agreement with those to be expected for samples not subject to deep cracking. Most of the deep cracks previously formed were 0. 045-inch or less (Table 11). The removal of 0. 046-inch apparently removed all cracks. Since a thermocouple is a joint between chromel and alumel, even when only chromel wire was used to attach the thermocouples there was a good chance of some contact with alumel from the thermocouple itself. This was manifested in a reduced tendency for deep cracking (Spec. R.-196 and Spec. R.-189 -- compared with Spec. R-193). However, alumel wire fastening or thermocouple contact itself did not completely explain the cracking. Specimens which had exhibited deep cracking in tensile tests were cooled to -320~F and broken at other locations by striking. In no case were other deep, oxidized cracks found than the one in the tensile test fracture. This suggests that the cracking started at one point and proceeded rapidly before it had a chance to occur at other points. Apparently deep cracking did not occur until creep strain exceeded 4 percent at 1200~ and 1300~F. It did not occur at higher temperatures. The presence of only one crack together with no apparent effect on creep curves up to 4-percent strain (Fig. 57) indicates that alumel catalyzed the cracking at one particular location and the crack grew quite rapidly. Otherwise, there should have been noticeable acceleration of creep and cracks present at the location of the other thermocouples. The absence at higher temperatures seems to be evidence that the reaction with alumel was not specific to the development of cracks at these temperatures. Examination of As-Exposed Surfaces Microstructural examination of as-exposed surfaces was carried out. The results showed that depending on the exposure temperature there was normally a relatively shallow zone of intergranular cracks and an area of oxidation together with probable removal of alloying elements by diffusion to the surface. In addition, the deep cracks previously discussed as arising from contact with alumel wire were observed in specimens exposed to more than 4-percent creep at 1200~ and 1300~F. Cracking Characteristics No cracking or other attack was noticeable (Fig. 58) in the specimens exposed for 100 hours at 1000~F without stress. Exposure at 1200~ and 37

1300~F under stress resulted in intergranular cracks (Figs. 59 and 60), usually one to a few grains deep. The depth of these surface cracks was less than 0. 001 to 0. 005-inch for specimens exposed to 2 to 4 percent creep (Table 11). The crack depth increased with the amount of prior strain (Fig. 54). The 0. 0125-inch of metal removed during re-machining removed these cracks, accounting for the absence of any effect on properties. When tensile tested with as-exposed surfaces, these cracks apparently were responsible for the reduced ductility. No cracks and no reduction in properties resulted from exposure without stress at 1200~ and 1300~F. When deep cracking occurred for creep strain of more than 4 percent due to contact with alumel, crack depths (Table 11) as deep as 0. 081-inch were observed. Measurements of the indicated depth of cracking by oxide coloration on specimens crept to rupture in 5. 1 to 9. 2 hours at 1300~F showed cracking depths ranging from 0. 07-to 0. 11-inch at the time of rupture. The deep cracks were intergranular. The reason why deep cracking occurred for more than 4-percent creep at 1200~ and 1300~F when in contact with alumel wire is not clear. These temperatures are those for minimum ductility and maximum notch sensitivity for the alloy. However, the cracks did not grow deeply in the absence of alumel. The brittleness at 1200~ and 1300~F alone could not be responsible. Some reaction with alumel causing a stress corrosiontype of effect at the grain boundaries was necessary. The notch sensitivity at these temperatures may have made the material more sensitive to attack by alumel and could thus explain its lack of effect at higher temperatures. General Surface Alteration When exposure temperatures were increased above 1300~F, general surface alteration occurred (Figs. 61, 62, and 63). There was penetration into grains and not just at grain boundaries. In addition, there was a layer of alloy depletion. This was probably due to decarburization and demolybdenumization as well as depletion of Al and Ti to cause the disappearance of gamma prime. Some internal oxidation and nitrogenation may have also occurred as is more evident in specimens exposed for prolonged times(Figs. 64 and 66). Measurements of the depth of visible general surface attack were made (Table 12) and plotted as a function of exposure time in Figure 68. The curves agree in general with a parabolic rate law similar to scaling (Ref. 19). The depth of detectable attack corresponded with the amount of attack first reducing properties. 38

The visible depth of penetration was correlated with ductility measurements (Fig. 69). This figure shows, as a function of depth of penetration, the ratio of ductility for specimens in the as-exposed condition to ductility with the surface attack machined off. The maximum effect was obtained with relatively little penetration. There was little difference between 10 and 100 hours of exposure. Evidently, only about 0. 0002-inch of penetration was required to give a maximum effect. Beyond this amount, there was no further deterioration. This was the cause for the absence of any change from 10 to 100 hours exposure. This also shows why there was relatively little additional effect from exposure above 1400~ to 1600~F Figure 69 indicated that Specimen R-188 (Table 9) should have had a visible attack depth of 0. 0003-inch after 10 hours exposure at 1600~F to 2. 81-percent creep. The surface was remachined to remove 0. 001inch and then the specimen tensile tested. The results (Fig. 55), however, showed that the properties were still the same as for specimens with asexposed surfaces. Examination showed that all of the surface attack had not been removed. Evidently creep did accelerate attack but it was not evident in the tensile properties due to the saturation at about 0. 0002-inch penetration. Structural Characteristics of General Surface Alteration The details of the microstructure of the altered surfaces are shown by Figures 70 and 72 for samples exposed 474 and 2012 hours at 1700~F without stress. The overaged gamma prime of the matrix is bounded by an area free from gamma prime. Figure 70 shows by the arrow that there was little, if any, preferential penetration along grain boundaries. The carbides and massive gamma prime are, however, removed from the grain boundaries in the layer next to unaltered matrix. The electron micrograph of Figure 71 shows in greater detail that no gamma prime was present. The next layers contained gamma prime-free matrix and a new phase in the form of needles. There was a layer next to this with some oxidation fissures as well as needles. The outside layer was free from needles but contained numerous fissures. Electron micrographs of the unattacked interiors of these specimens are shown in Figures 65 and 67. Four layers were machined from the surface of Specimen D-2 which had been exposed for 472 hours at 1700~F without stress. These layers were subjected to x-ray diffraction analysis. There was a linear decrease in the matrix lattice parameter (Fig. 73) from the base material out through the altered layers. Comparative parameter data (Fig. 74) from the work of Beattie and Ver Snyder (Ref. 20) on the effect of molybdenum on lattice parameters of a somewhat similar alloy make a good case for 39

the change in parameter being due to de-molybdenumization. There was good agreement with the present material for the lattice parameter at the Mo content of R.ene' 41 and also at the parameter of 3. 564A for the Mofree material and the surface of the present material. It does not seem possible that the lack of Mo would prevent gamma prime from forming. It seems quite likely, however, that Al and/or Ti were preferentially removed from solution in the alloy to account for the absence of gamma prime. Extraction residues from the layers machined from Sample D-2 showed considerable amounts of TiN (Table 13) in the altered surface. This points to the needle-like structure at least being TiN. This, in itself, could account for the absence of gamma prime. The data also suggest that A1203 was concentrated in the scale (Ref. 21). The presence of TiN plus A1203 probably account for the removal of Ti plus Al from solution. Both were, therefore, removed from gamma prime (i. e., Ni3(Al, Ti) ) to form more stable compounds. Even a partial removal of Ti plus Al could result in gamma prime depletion by lowering its solution temperature. The outside scale was mainly Cr203 as would be expected. It was also present in the sub-scale layer as the microstructure indicated. There was also some possibility of NiCr204 (Ref. 22). The layer next to the matrix showed M6C and M23C6, although these could have come from inadvertently including some of the unaltered matrix in the layer. The absence of carbides in the outer layers would be expected from the exposure conditions being favorable for decarburization. GENERAL PRINCIPLES FOR CREEP-DAMAGE OF GAMMA PRIMESTRENGTHENED ALLOYS The results of this investigation permit definition of several general principles regarding damage from elevated temperature creep-exposure to mechanical properties of nickel-base Ti+Al alloys strengthened by precipitation of gamma prime phase. These can be grouped into two broad categories. The first group includes principles for effects which can be considered reversible, that is, recoverable by heat treatment: 1. The measured volume fraction of gamma prime in the matrix decreases with exposure temperature and time. Yield and tensile strengths will decrease mainly as a result of the decrease in volume fraction of gamma 40

prime and only secondarily as a function of the accompanying increase in particle size. When the measured volume fraction reaches a constant value, the yield strength remains constant even though the particle size continues to increase. For Rene' 41, the measured volume fraction of gamma prime decreased during exposure above 1400~F. It tended to reach constant values in 100 to 300 hours at 1600~ to 1800~F. 2. The thermally-induced formation of massive carbides at the grain boundaries reduces ductility. These apparently can be both M23C6 and M6C with the total amount being the controlling factor. For Rene' 41, ductility began to be reduced after creepexposure for 10 hours at 1600~F and continued to decrease with higher exposure temperature and longer times. 3. Creep, in itself, had very little effect on mechanical properties at room temperature in comparison to the thermally-induced effects. Its main effect was for exposures below 1600~F and was to increase tensile yield strength and decrease compressive yield strength through residual stresses of the Bauschinger-type and to increase ultimate strength and decrease ductility through strain hardening. 4. The changes in mechanical properties from thermallyinduced structural changes could be restored by heat treatment. Restoring the gamma prime dispersion by re-heat treatment restored yield strength but not ductility. Restoration of ductility requires the re-solution of carbides as well. The restoration of gamma prime dispersion nearly restores creep-rupture strength. Re-solution of carbides as well as gamma prime is required for complete restoration. Recovery effects during re-heat treatment remove the Bauschinger effect and relieve strain hardening. The second group includes principles for those effects which cannot be recovered by heat treatment and are in this sense irreversible: 5. The lack of damage from creep itself and the ability to restore properties by heat treatment after creep is 41

probably unique to Rene' 41 due to its freedom from microcracking. Most other alloys of the nickel-base gamma prime-strengthened type are subject to microcracking during creep and, therefore, probably creep damage as well as thermally-induced structural damage. Microcracking would be expected to reduce all subsequent properties both due to a decrease in the loadcarrying area and the possible introduction of stressconcentration effects. The deterioration of properties by thermally-induced structural changes in Rene' 41 is a major factor in the reason for lack of creep damage. Apparently, deterioration of strength is so much greater than creep itself that very few of the expected manifestations of creep damage were observed. 6. Surface damage during creep exposure also influences properties when it is not removed by re-machining, Creep caused reduction in ductility after exposures as low as at 1200~F by inducing intergranular cracking at the surface. There was evidence that contact with alumel accelerated surface cracking. At 1400~F and above, the surface was damaged by other reactions as well as cracking. However, the damage from exposure without creep was as much as when creep occurred. Moreover, it reached its maximum effect in less than 10 hours even though the visible surface damage increased markedly with exposure time. The temperature of exposure was not a major factor since there was no further effect once the minimum attack requirement was attained. Surface damage did not influence yield strength. When the reduction in ductility was sufficient from either surface attack or structural changes, the tensile strength was reduced unless the yield strength was very high. With the exception of the absence of microcracking and, therefore, the absence of permanent damage, it is believed that these general principles developed from studies of Rene' 41 have wide applicability to all nickel-base Ti+Al strengthened alloys. Presumably, there will be differences in the details of the effects with composition variations. The general trends, however, should otherwise be the same. It is important, however, to recognize that the structure of the alloy and the marked influence of the instability of gamma prime are rather unique to alloys of this type. Other types of alloys might be quite different in response to creep damage. 42

CONCLUSIONS Mechanical properties of Rene' 41 alloy were changed by creep exposure through these main mechanisms: 1. Yield strength was reduced by a thermally-induced decrease in the measured volume fraction and, to a lesser extent, by an increase in particle size of the strengthening precipitate gamma prime in the matrix after exposure at temperatures above 1400~F. 2. Ductility was reduced progressively by a thermallyinduced increase in massive carbides in the grain boundaries with increasing temperature and time of exposure at 1400~F and higher. When the carbide build up was sufficiently severe, this also reduced ultimate tensile strength but not yield strength. 3. Tensile creep mainly increased tensile yield strength and reduced compressive yield strength by a Bauschinger effect when the exposure temperature was below about 1500~F. For such creep-exposure temperatures, recovery from the responsible residual stresses was insufficient to remove the effect. Some strain hardening also occurred at these temperatures. Creep had relatively little effect on the gamma prime effects. 4. Thermally-induced surface reactions reduced ductility over and above that due to carbides for exposure temperatures above 1400~F with apparently no added effect from creep. The maximum extent of reduction in ductility was attained in a very short time and, once attained, was independent of the exposure temperature. These reactions included loss of molybdenum, oxidation, and dec arburization. 5. Creep progressively reduced ductility at 1200~ to 1400~F by inducing surface cracking. There was evidence that contact with alumel accelerated the surface cracking. The thermally-induced structural changes were not permanent. After exposure, restoration of the volume fraction and particle size of gamma prime by re-solution and re-heat treatment restored yield strength. When the temperature of re-solution was high enough to dissolve M6C carbides, 43

both strength and ductility could be restored. Creep-rupture life was nearly restored by re-dissolving and re-precipitating gamma prime and was completely restored by re-solution of both gamma prime and carbides. The surface damage was not restored by re-heat treatment. No evidence of creep-induced microcracking was found in Rene' 41 alloy. Because this is not typical of many nickel-base Ti+Al hardened alloys, the lack of damage from creep itself may not be typical. The pronounced weakening from overaging of gamma prime may also have been a factor in the lack of permanent damage. The thermally-induced structural changes were accompanied by a volume decrease. The data suggest that this was due to removal of oddsized atoms from solid solution by precipitation reactions. The yield strength decreased to a minimum in 100 to 200 hours at 1400 to 1800 F which apparently coincided with attainment of near minimum measured volume fraction of gamma prime in the matrix. The particle size of undissolved gamma prime continued to increase with exposure time. The volume fraction of gamma prime was the controlling factor; an observation in accordance with published theory for dispersion strengthening. The results are believed to be reasonably typical for nickel-base Ti+Al hardened alloys, except for the absence of microcracking. A series of general principles were formulated for damage to mechanical properties of such alloys. There are, however, many points about the results, particularly the mechanisms of the effects, where additional research would be desirable for better verification. 44

REFERENCES 1. Gluck, J. V. and Freeman, J. W. "Effect of Creep-Exposure on Mechanical Properties of Rene' 41", ASD TR 61-73 (1961) 2. Gluck, J. V., Voorhees, H. R., and Freeman, J. W. "Effect of Prior Creep on Mechanical Properties of Aircraft Structural Metals", WADC TR. 57-150, Parts I, II, III. Part I: 2024-T86 Aluminum (1957), Part II: 17-7PH (TH 1050 Condition) (1957), Part III: Cl1OM Titanium Alloy (1958) 3. Gluck, J. V. and Freeman, J. W. "Effect of Prior Creep on ShortTime Mechanical Properties of 17-7PH Stainless Steel (RH 950 Condition Compared to TH 1050 Condition", WADC TR 59-339 (1959) 4. Gluck, J. V. and Freeman, J. W. "Effect of Prior Creep on the Mechanical Properties of a High-Strength Heat-Treatable Titanium Alloy, Ti-16V-2.5A1", WADC TR. 59-454 (1959) 5. Gluck, J. V. and Freeman, J. W. "Further Investigations of the Effect of Prior Creep on Mechanical Properties of Cl1OM Titanium with Emphasis on the Bauschinger Effect", WADC TR 59-681 (1960) 6. Letter from W. H. Couts, Jet Engine Dept., General Electric Co., Dated January 7, 1960 7. Bigelow, W. C., Amy, J. A., and Brockway, L. O. "Electron Microscope Identification of the Gamma Prime Phase of NickelBase Alloys", Proc. ASTM, v. 56, p. 945 (1956) 8. Nelson, J. B. and Riley, D. P. Proc. Phys. Soc. v. 57, p. 160 (1945) 9. Gluck, J. V. and Freeman, J. W. "A Study of Creep of Titanium and Two of its Alloys", WADC TR 54-54, p. 83 (March 1956) 10. Fountain, R. W. and Korchynsky, M. "The Phenomenon of'Negative Creep' in Alloys", Trans. ASM, v. 51, p. 108-122 (1959) 11. Beattie, H. J. "Aging Reactions in Rene' 41", General Electric Co. Report No. DF59SL314 (May 1959) 45

12. Morris, R. J. "Rene' 41... New Higher Strength Nickel Base Alloy", Metal Progress, v. 76, No. 6, pp. 67-70 (December 1959) 13. Rowe, J. P. "Relations Between Microstructure and Creep-Rupture Properties of Nickel-Base Alloys as Revealed by Over-Temperature Exposures", University of Michigan PH. D. Thesis, 1960 14. Bunshah, R. F. and Goetzel, C. G. "A Survey of Dispersion Strengthening of Metals and Alloys", WADC TR 59-414 (March 1960) 15. Meiklejohn, W. H. and Skoda, R. E. "Dispersion Hardening", Acta Metallurgica, v. 8, No. 11, p. 773 (November 1960) 16. Dew-Hughes, D. "On Meiklejohn and Skoda's letter'Dispersion Hardening' ", Letter to the Editor, Acta Metallurgica, v. 8, No. 11, p. 816 (November 1960) 17. Bigelow, W. C. and Amy, J. A. "Electron Metallographic Studies of Nickel-Base Heat-Resistant Alloys", WADC TR 58-406, ASTIA Document AD 155767 (August 1958) 18. Beattie, H. J. and Hagel, W. C. "Intragranular Precipitation of Intermetallic Compounds in Complex Austenitic Alloys", Trans. Met. Soc. AIME, v. 221, No. 1, February 1961, p. 31 19. Barrett, C. A., Evans, E. B., and Baldwin, W. M. "Thermodynamics and Kinetics of Metals and Alloys. High Temperature Scaling of Ni-Cr, Fe-Cr, Cu-Cr, and Cu-Mn Alloys", U. S. Office of Ordnance Research Tech. Report, Dec. 1955 (declassified) Summarized in The Nickel Bulletin, v. 32, No. 2 (February 1959), The Mond Nickel Co., London. 20. Beattie, H. J. and VerSnyder, F. L. "The Influence of Molybdenum on the Phase Relationships of a High Temperature Alloy", Trans. ASM, v. 49, p. 894 (1957) 21. Radavich, J. F. and Wilson, J. E. "Phase Identification in NickelBase Alloys", Metal Progress, v. 79, No. 5, p. 94 (May 1961) 22. Ignatov, D. V. and Shamgunova, R. D. "Mechanism of the Oxidation of Nickel and Chromium Alloys", NASA Technical Translation F-59, March 1961 (published in Russian, 1960) 46

Table 1 Effect of Long-Time Aging on Tensile Properties of Rene' 41 Exposure Conditions Room Temperature Tensile Properties After Exposure Spec. Temp Time Ult, Tensile Str,.2% Offset Yield Str. Elong, Red, of Area Modulus, E No. (~F) (hr) (psi) (psi) (%) (%) xl06 psi As Treated 189,800 129,733 20.5 27.9 30,5 R-59 1600 200 172,100 106,000 13.5 15.0 31.3 D-3 1600 401 142,400 91,500 7,6 11.0 30.0 R-76 1700 200 128,000 96,400 3,8 6.0 32.2 D-2 1700 474 137,200** 101,800 3.3 3.3 31 8 D-5 1700 2012 141,400 94,500 6, 0 8. 0 380 R-78 1800 200 155,000 109,800 6.3 8,5 32.2 D-4 1800 1150 142,000 83,000 6.0 9.7 34 5 D-6 1800 1700 149,000 -- 5.2 6.7 31 1 Notes: D = dilatometer aging R = conventional "unstressed" exposure (54. 5 psi) * = "R" condition: 1950~F - 1/2 hr + AC; 1400~F - 16 hr + AC ** = Broke at gage mark

Table 2 Impact Test Data for Rene 41 Tested at Room Temperature Exposure Conditions Impact Properties Temp Time Stress Type of Smooth or Impact Strength ( ~F) (hrs) Test * Notched ft-lb As Heat Treated Izod N 3 Izod N 3 avg. 3 Charpy N 3.5 Izod S >30 ^ >30 avg. >30 Charpy S 37 17 34 avg. 29 600o 10 none Charpy S 29 Charpy S 27 avg. 28 1600 401 none Charpy S 6 1700 10 none Charpy S 16 Charpy S 27 avg. 21,5 1800 10 none Charpy S 18 Charpy S 12 avg. 15 Izod Test - Specimen Charpy Test - Specimen Supported as Vertical Cantilever Supported as Horizontal Beam Specimen Dimensions: 2. 16 x o 197 x. 197 inches If Notched: 45~V notch at mid-point 0O 039" deep - 0o 010 root radius.^... Specimen bent but did not break 48

Table 3 Establishment of Re-Heat Treatments for Restoration of Room Temperature Tensile Properties of Rene' 41 Room Temperature Tensile Properties After Exposure, Initial Exposure Conditions Re-Heat Treatment (if any), and Re-machining _ Heat Temp* Time Stress Total Load Plastic Load Creep Def. Total Def. Re-Heat Code Ult. Tensile.2% Offset Elongation Red, of Area Spec. No. Treatment (IF) (hrs) (psi) Def. (%) Def. (%) (%) (%) Treatment (if any) Str. (psi) Yield Str.(psi) (%) (%) Mill Anneal (As Received) -- -- -- -- -- -- -- 128, 400 57, 400 52. 0 57.9 T90F- 72 hr+AC --- ------ -- - -- -- R 1400~F - l6hr+AC -- -- -- -- -- -- avg. 189,800 129,733 20.5 27.9 Exposure: R-32 R 1800 100,.0 Nil -- -- -- -- None -- 137,500 102,600 4.2 5,9 R-104 R 1800 100,.0 Nil -- -- -- -- None -- 149,800 107,200 6.7 7,2 avg. 143,650 104,900 5.4 47 Exposure Plus Re-Solution of Gamma Prime and M23C6 Carbides Only: R-136 R 1800 100,.0 Nil -- -- -- -- 1950~F - 1/2 hr+AC) 1400~F - l6hr+AC] R 160,200 121,800 6.2 6,8 Exposure Plus Complete Re-Solution and Re-Heat Treatment: R-105 R 1800 100,.0 Nil -- -- -- -- 2150F - I hr+AC 1975'F - I hr+WQ 1950~F - 1/2 hr+AC 1400~F - l6hr+AC -- 179,800 121,000 17,6 1862 R-110 R 1800 100,.0 Nil -- -- -- -- 2150F - 2 hr+AC 1975'F - I hr+WQ 1950~F - 1/2 hr+AC 1400~F - 16 hr+AC -- 183,000 122,000 21.1 21,4 R-1 1 R 1800 100.0 Nil -- -- -- -- 2150F - 2 hr+AC 19751F - I hr+WQ A _- - - - -___ __ __ __ __ _ _ __ __ _ _ _ _ _ _ _ _ 1400F - 16 hr+AC J _ 179, 200 119,000 23.4 25., - R-201 R 1800 100,.0 Nil -- -- -- -- see above A 183,200 127,200 24.0 28,2 R-135 R 1800 100,.0 Nil -- -- - -- -- 2150F - 2 hr+AC 1400F - l6hr+AC 185,000 126,000 24.0 256 Complete Re-Solution Without Exposure: R-141 R none --- -- -- -- -- -- see above A 174,000 111,000 30.4 36,3 Complete Re-Solution After Exposure to Creep at 1800~F: R-26 R 1800 100.0 7500 0.04 Nil 1.70 1.74 None -- 130,200 106,800 3.4 11.7 R-45 __ R 1800 100.0 7900 0.07 _ Nil 5. 17 _ 5.24 None -- 139,500 119,000 3.5 3 4 R-124 R 1800 100.0 7000 0.04 Nil 3.46 3.50 see above A 173,200 115,200 22.6 31,2 R-116 R 1800 100.0 7500 0.06 Nil 28.40 28.46 see above A 170,500 116,000 16.9 2092 Re-Heat Treatment of Material Originally Heat Treated at 2050' and 1650'F: 2050~F - 1/2 hr+AC 1650~F - 4 hr+AC -- -- -- -- -- -- R2 180,300 115,100 25.3 29.6 R2-8 R2 1800 1000 Nil ___ -- -- -- - None ___ — 140,500 105,100 5.0 4.9 R2-13 R2 1800 100.0 Nil -- -- -- -- 2050~F - 1/2 hr+AC 1~ -_ -__ ___ -_ -_ -_ -_ -_ -__ - - -___ -_ -__ -__ -__ ___ -_ -__ -_ -_ ___1650~F - 4 hr+ACJ J 169,500 113,800 11. 3 130 R2-14 R2 1800 100.0 Nil -- -- -- -- 2150F' - 2 hr+AC 2050~F - 1/2 hr+AC 1650~F - 4 hr+AC 176,200 113,400 19.9 20.8 R2-15 R2 1800 100.0 Nil -- -- -- -- 2150F- 2 hr+ACT A 1650F - 4 hr+ACJ 177,200 115,000 19.8 21.4 * plus 4 hours pre-heat

Table 4 Effect of Re-Heat Treatment on Tensile Properties of Rene' 41 After Initial Creep-Exposure Room Temperature Tensile Properties After Re-Heat Treatment and Exposure Conditions Refmachining. 025-inch from Dia., Temp Time** Otress Frac, of Rupt. Total Load Plastic Load Creep Def. Total Plastic Total Def. Re-Heat. Ult. Tensile.2% Offset Elongation Red. of Area Modulus, E Spec. No. (~F) (hrs) (psi) Life (est.) Def. (%) Def. (%) (%) Def. (%) (%) Treatment* Str. (psi) Yield Str. (psi) (%) (%) xl06 psi As Treated "R" - - - - - - - - - 189, 800 129, 733 20. 5 27.9 30.5 R-150 1400 10.0 84,700 0.91 0.38 Nil 3.14 3.52 3.52 R 182,800 116,000 23.5 31.4 30.6 R-172 1400 100.0 61,000 0.80 0.24 0.01 19.0 19.0 19.01 R 192,200 132,000 19.6 38.0 29.4 R-148 1600 10.0 39,000 0.71 0.21 0.02 6.20 6.20 6.22 R 185,500 118,800 22.4 29.6 3168 R-185 1600 100.0 23,500 0.74 0.12 Nil 7.38 7.38 7.50 R 182,200 137,200 12.0 13.6 31.4 R-171 1600 100.0 23,500 0.74 0.10 Nil 25.50 25.50 25.60 R 183,000 132,000 10.3 12.5 30.1 U-I 0 R-151 1700 10.0 23,000 0.60 0.12 0.01 3.56 3.57 3.68 R 190,000 142,500 15.6 17.0 33.8 R-155 1800 9.2 12,600 0.65 0.07 Nil 5.34 5.34 5.41 R 187,000 138,800 13.5 15.2 29.1 R-197 1800 85.5 7,900 0.74 0.04 Nil 8.06 8.06 8.10 R 161,400 140,800 3.1 5.2 30.9 R-141 As Treated "R' - - - -- - - A 174,000 111,000 30.4 36.3 3134 R-116 1800 100.0 7,500 0.77 0.06 Nil 28.40 28.40 28.46 A 170,500 116,000 16.9 20.2 29.8 R-124 1800 100.0 7,000 0.59 0.04 Nil 3.46 3.46 3.50 A 173,200 115,200 22.6 31.2 32,2 As Treated"R2" -R- - - - - - --- 180,300 115,100 25.3 29.6 31,2 R2-16 1800 100.0 7,500 0.77 0.03 Nil 1.54 1.54 1.57 A2 168,200 111,800 15.4 16.5 31,6 /R 1950' - 1/2 hr + AC; 1400~ - 16 hr + AC R2: 2050~ - 1/Z hr + AC; 1650~ - 4 hr + AC A: 2150~ - 2 hr + AC; 1975~ - 1 hr + WQ; 1400 - 16 hr + AC A2: 2150~ - 2 hr + AC; 1650~ - 4 hr + AC ** Plus 4 hours pre-heat prior to load application

Table 5 Effect of Re-Heat Treatment on Creep-Rupture Properties of Rene' 41 After Initial Creep Exposure Expected Life Re eat treartd pluSj R oracio First Exposure Remaining Second Exposure remacined Recovered o Temp Stress Time** Normal Life Frac. Creep Def. Red. of (Uninterrupted Est. Rupt. Rapt. Time** Elong. Red. of Total Liie Lfe nd Exo Fraction Recovered Total Life Total Total Eln. pended Spec. No. (*F) (psi) (hr.) Rupt. Life (hr.) Expended (%) Area(%) Test) (hrs) Elong. (%) (hrs () rea(% ) (hrs) Exc td. Life Lnd Exp. Of Eended Lie Normal Lie Elong. Est. Rupt Elo uctlt Original Condition: "R": Re-heat Treatment "R"* R-153 1400 84.700 10.0 11.5 0.91 4.29 5.7 1.5 25 6.9 a) 3,7 7.6 17.6 4.60 054.53 8.0 0.32 -3 97 R-179 1400 61.000 100.0 125.0 0.80 14.5 16.4 25.0 25 85.9 18.3 34.3 185.9 3.44 0 61 1.49 32.8 l31 0.54 R-149 1600 39 000 10.0 14. 0 0.7 1 7. 3 7.8 4. 0 25 _ 14. 0_.9 __ 41.5 24.0 3.50. 0 _ __ 0.- - _l-7rOO —*T^ T- 5 - 50 -- 25 b) 11.1 213.5 -2.4 2 -— 1.- -..-78 0.71 1. 15.9 0. -3.30 R-183 1600 39.000 10.0 14..0 05 R-187 1600 39,000. 10,0 14.0 0,71 - 2.73 _3.0 4.0 25 c) 8.4 2.25 43.4 18.4 20 0.44 131 25 Z 1 R-178 1600 23,500 100,0 135.0 0.74 10.76 1.02 35.0 26 85.7 26.4 41.5 185.7 2.45 0.51 1.37 371 1.43 1.04 R-154 1700 23,000 9.0 15.0 0.60 6.28 8.8 6.0 27 10.9 28.4 52.3 19.9 1.82 0.55. 32 34.7 1.,9 1.3 (T R8-157 1800 12,600 6.8 10.4 0.65 5.34 6.2 3.6 34 8.6 33.8 51.8 15.4 2.39 0. 73 1.48 39.1 1.15 0.96 R-181 1800 7,300 100.0 145.0 0.69 8.36 9.0 45.0 41 94.0 33.0 46.0 194.0 2.09 0 49 1__34 41.4 1.01 Original Condition: "R": Re-heat Treatment "A"* R-126 1800 12,500 10.0 12,0 0.83 7.77 8.8 Z.0 34, 23.4 20. 0 42,.0 33.4 11.70 2. 14 2.78 Z7.8 0.82 -0.80 R-132 1800 9,500 33.0 45.0 0.73 9 10 11.0 12.0 43 93.9 18.6 46.6 126.9 7.83. 48 2.82 27 7 0.64 -1.68 R-133 1800 7,300 100.5 145.0 0.69 6.93 7.3 45.0 43 414.5 10.6 33.1 515.0 9.20 3.70 3. 54 17.5 0.41 -3.68 R-166 1800 1Z,500, - - - - - - 21.4 32.0 49.0 - - - - R-167 1800 9.500 No flrst - - - - - - 67.5 34.0 53.6 - - - - R-173 1800 9500 - - - - - - 144.1 2.0 36.0 - - - - - R-168 1800 7:300 Eoposure -.. - - 357.8 20.6 32.8 - -... - Original Condition: "R": Re-heat Treatment "R3"* R-170 1600 23,500 100.0 135.0 0.74 3.78 3.9 35.0 26 122.0 19.8 35.2 222.0 3.49 0.87 1.64 23.6 0,95 -0.62 Original Condition: As Rec. (Mill Anneal) Heat Treatment "A"* R-174 1800 9,500 Nofirst exposure - 126.1 18.8 30.09'9 - * Heat Treatments "R": 1950 - 1/ hr + AC "A": 2150 - hr + AC "R3": 1950 - 1/2 hr AC ) premature failure? b) Re-heat treated - NOT remachined 1400- 16 hr + AC 1975 - L hr + WQ 1600 - 1/3 hr + AC ) Remachined ONLY 1400 - 16 hr + AC 00 PluC 4 hours pre-heat prior to load application

Table 6 Structural Parameters and Yield Strengths of Rene' 41 after Unstressed Exposure Average Average Meiklejohn-Skoda Exposure Exposure Inter-Particle Particle Parameter Temp Time Volume Fraction Spacing,X Size, h f 0. 2% Offset Spec. No. (~F) (hrs) of y' (f) ____ ____ A 0.82 - fr Yield Str., psi As Treated - -.37 546 1000 7.03 130,000 R-52 1600 14.27 512 870 3.72 129,000 R-31 1600 104.23 835 1245 2.68 109,000 R-59 1600 204.22 1110 1610 2.66 106,000 D-3 1600 401.18 2250 2780 2.22 91,500 (i R-42 1700 14.25 1010 1600 3.31 108,200 N R-73 1700 104.21 1945 2660 2.64 94,500 R-76 1700 204.18 2520 3110 2.22 96,400 D-2 1700 474.21 2540 3480 2.64 101,800 D-5 1700 2012.19 3155 3990 2.35 94,500 R-61 1800 14.24 1635 2520 3.12 102,500 R-32* 1800 104.19 1025 1300 2.35 102,600 D-4 1800 1150.12 5120 4650 1.52 83,000 R-2 As HT.26 755 1240 3.50 115, 100 R2-7 1600 104.21 1035 1440 2.64 108,000 R2-12 1800 14.14 1945 1945 1.73 98,300 * fine y' reprecipitated on cooling

Table 7 Comparison Matrix Lattice Parameters and Dilatometer Shrinkage for Rene' 41 Aged Without Stress Spec. Aging Lattice Calculated Dilatometer No.* Treatment Parameter Shrinkage Shrinkage As Heat Treated - 3. 6050 D-3 1600~F - 401 hr, 3, 5996 0.15% 0,18% D-2 1700~F - 474 hr. 3, 5948 0.,29% 0,21% Un D-5 1700~F - 2012 hr, 3. 5940 0,31% 0.,36% D-4 1800~F - 1150 hr. 3,5926 0.,34% 0.40% * All Specimens Originally in Condition "R": 1950~F - 1/2 hr + AC 1400~F - 16 hr + AC Comparative Results from Data of Beattie 2150~F - 2 hr + WQ 3. 607 (Ref. 11) above plus 1600~F - 1000 hr. 3. 592 0.,42%1

Table 8 X-Ray Diffraction Data from Extraction Residues of Rene' 41 Specimens Used for Development of Re-Heat Treatments Sp ec. No. R-104 R-105 _ __ __ R-lll|11 Standard Patterns for Indicated Phases Exp. Temp. Time 1800~ - 100 hr 1800' - 100 hr 1800~ - 100 hr Re-Heat Treatment None 2150~ - 1 hr+AC 2150' - 2 hr+AC 1975~ - 1 hr+WQ 1975~ - 1 hr+WQ 1950" - 1/2 hr+AC 1400~ - 16 hr+AC 1400~ - 16 hr+AC 6 23 6 Tensile Test Elong. 6.7 17.6 23.4 a = 4.32 a = 11.10 a = 10.71 Red. of Area 7.2 18.2 25.7 "d" Range "d" I "d" I "d" I "d" I "d" I "d" I ___.______ 2.89 5 2.86 20 2.88 2.70-2.79 2.76 5 2.77 5 2.77 m 2.60-2.69 2.67 2.68 w 2.50-2.59 _ ____2. 54 ms 2.39-2.49 2.48 80 2.49 100 2.49 100 2.49 s 2.48 2.41 5 2.46 vw 2.40 40 2.30-2.39 2_. 36 5 2. 38 10 2.. 399 s 2.20-2.29 2.26 30 2.25 10 2.25 10 2.263 s 2. 10-2. 19 2. 19 40 2.185 s 2.15 60 2.15 100 2.15 90 2.16 s U1P^ ___________.2. 13 60 2. 13 10 l 2. 135 s 2.00-2.09 2.065 100 2.067 10 2.06 80 2.062 s 2.03 5 2.00 40 1.89-1.99 1.96 20 1.96 5 1.959 m 1.80-1.88 1.88 30 1.875 w 1.892 m 1.84 30 1.85 5 1.850 w 1.81 10 1.82 30 1.8 s 1.69-1.79 1.79 10 1.753 vw 1.785 mw 1. 71 5 1.74 20 1.69 w 1.692 vw 1.60-1.68 1.66 10 1.67 w 1.632 vw 1.62 10 1.615 ms 1. 50-1.59 1.55 5 1.552 mw 1.545 vw 1.52 50 1.53 80 1.52 60 1.528 m 1.528 w 1.500 vw 1.40-1.49 1.43 5 1.481 1.431 1.40 5 1. 443 mw 1.30-1.39 1.35 5 1.385 1.395 vw 1.303 70 1.30 50 1.302 mw 1.353 1.340 w 1.27 20 1. 15 40 1.03 30 TiC TiC TiC M6C medium M6C slight M6C M23C6 medium M23C6 slight-medium M23C6 "d" - interplanar spacing I - relative intensity (to strongest line)

Table 9 Effect of Remachining On Response of Room Tempereture Tensile Properties of Rene' 41 To Creep-Exposere Exposure Conditions Room Temperature Tensile Properties After Exposure % of Rupture Amount Ult. Tensile.2% Offset Temp Time** Stres Life Total Load Plastic Load Creep Def. Total Plastic Total Def. Remachlned Strength Yield Strength Elongation Red. of Area Mod.lu, E Spec. No.* (*F) (hrs) (psi) (eat) De. (%) De. ) () De. () ) from dia. (.) (psi) (psi) (i) (() (O06 psi As Treated - - - - - 189,800 129.733 20.5 27.9 30.5 5-89 1000 100.0 None Nil. - - - - 0.025 187.000 130,000 24.4 32.4 31.6 R-90 1000 100.0 None Nil - Nil 187500 129.500 22.9 23.2 31.2 R-146 1200 10.0 136,000 99 2.42 1.85 2.85 4.70 5.27 Nl 212.000 - 1.3 2.0 33.6 R-142 1200 10.0 134.000 90 1.50 1.00 1.26 2.26 2.76 Nil 202.500 189.500 14.2 17.6 32.4 R-128 1200 10.4 133,000 83 2.31 1.80 2.22 4.02 4.53 Nil 202,000 200.000 2. 6.1 30.0 R-112 1200 10.0 132,000 72 1.50 0.98 0.81 1.79 2.31 N1 197,500 181,000 13.1 18.6 30.2 R-o10 1200 10.0 130.000 57 2.7Z 1.90 1.77 3.67 4.49 Nil 203.000 192,.000 11.0 13. 2 29.8 R-101 1200 10.0 129.000 51 1.59 1.04 0.39 1.43 1.98 Nil 188,000 182.000 14.0 17.6 29.6 R.-100 100 120.0 None Nil - - 0.025 191,300 133,800 20.4 26.4 51.9 R-96 1200 10.0 None one Nil - - - - Nil 190.500 133.200 21.9 27.4 32.4 R-97 1200 100.0 None None il - - - - - Nil 194.500 138,000 21.4 25.5 31.6 Normal Thermocouple Procedure: R.140 1300 10.0 113,500 98 0.56 0.08 4.77 4.85 5.33 0.025 204.000 - 2. 1 3.0 R-127 1300 10.0 112.000 97 0.62 0.12 4.18 4.30 4.80 Nil 190,000 - 0.9 4.1 31. R-123 1300 10.0 111,000 91 0.62 0.15 2.96 3.11 3.58 0.025 200.000 181.000 16.9 29.9 31.7 R-102 1300 4.5 108.0004 70 0.70 0.20 1.70 1.90 2.40 Nil 198.000 166.300 15.5 14.8 31.3 R-95 1300 10.0 108.000 70 0.52 0.07 2.64 2.64 2.71 Nil 199.800 173.200 10.2 12.9 31.2 R-119 1300 10.0 108,000 70 0.50 0.05 0.63 0.68 1. 13 Nil 188.500 150.000 11.9 13.7 30.6 R -99 1300 10.0!06.000 66 0.49 0.04 0. 59 0.63 1.08 Nil 196,200 166,000 1 7. 6 17.2 30.4 R-134 1300 10.0 None Nil - - - - - - 191.800 133,800 21.6 24.9 31.4 No Thermocouples on Gage Section: R-145 1300 10.0 114.000 99 0.66 0.20 5.14 5.34 5.80 Nil 211,000 204.000 5.9 8.7 32.6 R-144 1300 10.0 114.000 99 0.53 0.12 1.78 1.90 2.31 0,025 197.000 166,800 15.6 27.8 31.2 R-130 1300 10.0 112,500 97 0. 55 0.09 3.65 3.74 4.20 Nil 204.000 188.000 8.0 11.3 32.4 R-115 1300 10.0 108.000 70 0.48 0.03 0.41 0.44 0.89 Nil 195,000 158.000 17. 3 19.2 31. 1 r One Thermocouple at Center of Gage Section + Single Loop. of Chromel and Alumel in Contact at Ends of Gage Section R-152 1300 9.2 114,000 99 0.64 0.23 5.36 5.59 6.00 Nil 208.000 203,000. 9 4. I. 0 H8all Gage Section Entirely Wrapped With Chromel, Half Entirely Wrapped With Alumel R-156 1300 8.7 114,000 99 0.60 0.22 5.66 5.08 6.26 0.091 210,000 207,000 6.2 22.6 29. 5 All Three Couples Tied on With Chromel x R-196 1300 8.4 113,500 98 0.68 0.23 4.41 4.64 5.09 Nil 208,000 200,000 6.4 6.8 30.1.B-189 1300 9.4 113,000 98 0.63 0.20 3.87 4.07 4.50 Nil 206,000 196.000 6. 5 6.7 30. 4 All Three Couples Tied on With Alumel -- 193 1300 7.7 113,000 98 0.70 0.24 3.80 4.04 4.50 Nil 196,500 192,500 2.0 4.1 31.6 riret Exposure -120 1300 10.0 108,400 83 0.52 0.05 0. 64 0.69 1. 1 Nil.. $Secod Exposure brought up to 8-*120 1300 90.0 86,000 75 temp. under load 4.50 4.0 4.. -eTlf 1300 T O -.0 Nil 186,000 14,.000 1.2 2.9 32, 6 - 147 1300 97.3 88,000 97 0.36 Nil 5.21 5.21 5. S7 Nil 193,000 185.000 f. 5 S. 33.2 R.143 1300 100.0 87.000 91 0.36 Nil 3. 11 3.11 3.47 0.025 199,800 166.400 17.4 03.4 30.0 81-113 1400 4.0 85.500 45 0.41 Nil 6.47 6.47 6.88 Nil 207,000 196.000 9. 6 14.8 38.4 8-109 1400 10.0 84.000 91 0.36 Nil 1.36 1.36 1. 72 Nil 196.800 153,000 20. 6 24.8 30.4 R-8 1400 10.0 80.000 57 0.36 Nil 2 (egt) 2( et) 2a (et) Nil 192.400 152,600 16.0 17.2 30.1 R-11 1400 10.0 None Nil... - Nil 183,500 134,000 14.7 18.7 29.7 R -9 1400 100.0 None Nil -.." - - Nil 181,000 128.000 14.9 14.9 31.6 lf13 1500 10.0 55.000 48 0.30 Nil 1.24 1.24 1.54 Nil 179.500 136.000 9.9 15.2 30.1 1.-125 1600 10. 0 39,000 72 0.20 Nil 6. 50 6. 50 6.70 Nil 184.000 147.500 9.1 10. 4 32.4 I-18~ 1600 10.0 39.000 72 0.17 Nil 2.81 2.81 2.98 0.002 178.000 138,200 9.8 9.8 31.2 I.9 1600 10.0 35.000 36 0.18 N11 I. 13 1. 13 1.31 Nil 169,000 128,200 8.9 11.5 30.2 8.10 1600 10.0 None Nil - - Nil 175, 800 125,500 10. 1 13.7 30.1 3.92 1600 100.0 None Nil.. - Nil 149,000 107,500 6.8 8.3 31, R-31 1800 10.0 12.000 77 0.07 Nil 4.45 4.45 4.52 Nil 135,000 111,000 3.5 5. 2 32.0 B* I2 1800 10.0 None Nil.. - Nil 158,000 104,300 9.5 10.5 29.9.-I11 1800 50.0 None Nil -. Nil 168.000 103,500 14.0 15.3 29.2 R-91 I800 100.0 None Nil. - - Nil 131.000 102.500 2.8 5.0 30.4 8 Data for remachined specimens through R-88 (from Ref. I) included in Table 10 - Plue 4 hour pre-heat prior to load appllcation Deviation from nornml thermocouple attachment procedure 55

Table 10 CREEP EXPOSURE TEST DATA FOR RENE' 41 ALLOY (REMACHINED SPECIMENS) of Temp Time* Stre. Rupture Total Load Plastic Load Creep De,. Total Plastic Total Def. Ult. Toertli 0. 25 Offet Elongation Reductlo. Modlu E Specimen No. (*) (hro) (pi) Life (eat) Def. (%) Def. (%) (%) Def. (%) (%) Strength(pai) Yield Str. (psi) () of Area( ) ~ 10 (pil) CONDITION "R": 1950'F - t/2 HR + AIR COOL; 1400'F - 16 HR + AIR COOL TENSION TESTS An Treated (avg) - - -- -- -- -- 189, 800 129.733 20. 27.9 30. R-66 1200 10.0 133. 000 80 2.94 1.86 1.63 3.49 4.57 200,000 185, 500 4.7 ) S. 3 31.9.-28 I1200 10.0 130. 00' 55 Z. 77 2.40 1.02 3.42 3.79 208.000 193.000 21.5 28.0 30.6 R -58 1200 100.0 114.000 91 0.54 0.08 0. 38 0.46 0.92 198.800 160.400 10.9 IZ. 0 32. 3 3t-34 1200 100.0 113.000 83 0.77 0.30 1.S2 1.82 2.04 199.000 161,000 14.9 25.0 30. R.74 1200 116.0 None Nil - -..- 193.200 133.000 21. 8 25.8 31.7 R-63 1300 10.0 108.397 83 0.50 Nil 4.60 4.60 5.10 171.000 -- 0.4) 0.7 33.4 1-38 1300 10.0 106.000 67 0. 0 0.04 0.0 4 0.46 0.50 0.89 191,000 143,000 Z0.1 26.4 32.2 R-21 1300 10.0 100. 000 34 0.49 Nil 0.31 0.31 0.80 203. 000 153,000 19.8 23.8 33.2 3.438 1300 100.0 86.500 87 0.40 Nil 1.24 1. 4 1.64 203. 000 158.000 19.3 23.2 31.3 R-33 1300 I00.0 85.000 77 0.39 Nil 0.83 0.83 1.22 201,000 157.200 20.0 27.1 31.8 a3-69 1400 11. 6 84.000 99+ 0. 35 0.02 2,95 2.97 3.20 194. 500 161.000 1S. 9 23.2 33. 3-44 1400 10. 0 83. 000 80 0. 35 Nil 2.69 2.69 3.04 203.000 172.000 16. 7 27. 7 31.2 3.-65 1400 10. 800.000 57 0.42 Nil 1.24 1.24 1.66 200.500 15S.000 18.5 28.4 32.6!:-S 1400 10.0 None Nil -195. 100 138, 300 20. 9 29.5 33.7 3.-S4.. 1400''20.5' 80.000' 99+ 0.36 Nil 7. 5 (et) 7. S(eot) 7.86 (eot) 207,000 181ZO O 8.6 27.6 30. 5' -40 1400'00. 61.000'' 0.45 Nl 4.40 4.40 4.8.20o.000 163..00 S 14. 5 22.6 31.6 R-23 1400 100.0 56. 000 48 0. 35 Nil 1,32 1.32 1.67 200,000 148.000 18.I 23.6 35.0 R-30 1400 100.0 None Nil -- -. - -- -- 190.000 131.000 22.0 24.9 31.1 R-36 1500 10.0 60,000 91 0.37 Nil 3.45 3.45 3.92 190.500 151.000 18.4 20.3 32.7 -55 1500 10.0 55.000 48 0.24 Nil 0.67 0.67 0.91 192.500 138.700 19.7 22.1 31.0.41' o1500 I100.0 39.500 80s 0. 17 0.08s 7, 3 (eat) 7. 35 (oet) 7.47 (eat) 192,000 145.000 12. 1 13.2 32.6 R -24 1500 100.0 35, 000 43 0. 20 Nil 0. 71 0. 1 0. 91 194.000 134.000 10.5 1.7 33. R-47 1600 10.0 39.800 80 0.18 Nil 8. 6 (et) 8.6 (otj 8.78 (oat) 185,000 143.400 10.0 13.6 30.6 3.43 1600 10. 34 800 38 014 Nil 1.07 1.07 1. 21 183.000 130.000 16.5 17. 30.4 3-52 1600 10.0 None Nil - - - - -- 190,500 129.000 20.4 25.9 30.8 - 71 1600 100.0 24.000 89 0 11 Nil 0.61 0.61 0.72 159.500 110.800 8.6 8.4 30.8 3-Z7 1600 t00.0 23,000 77 0. 13 Nil 1.35 1.35 1.48 170.000 114.100 10.3 11,.7 31.8 3-25 1600 100,0 19.000 45 0.09 Nil 0.21 0.21 0.30 171.000 111.400 10.8 13.7 32.2 3-31 1600 100.0 None Nil -- -- 0- o- 10.,000 106.000 15.1 16,6 30.1 -60 1600 200.0 19'.000 91 0.15 Nl 4.0 4.08 4,20 1 59 200 109,500 6.8 7.9 34.4 3.59 1600 200. 0 None Nil.. - — 172.100 106,000 13. 18.0 31.3 3o.37 170r 10.0 Z3. 00 71 00.13 NIl 14. t 1 () 14. 33 (t) 170.400 117,800 11.3 12.0 31.0 3.-70 1700 11.6 22,500 63 0,11 Nil 4,4 (oat) 4,4 (oat) 4.5t (oat) 167,500 115,300 10.3 11.4 31,9!-20 1700 10.0 21.000 48 0.1 0. 001 074 0. 0.36 1.5,000 118.000 13. 19.0 32.1 3-42 1700 10.0 None Nil 179,000 103,200 21 9 26. 3..6 36' 1700 100 " So12o,00 77 0 07'' Nl 3.04 3.04 3. 11 153,200 109,500.,2 I.0 33.2 -9 1710 100.0o 12.000 69 0. o0 NI 0.84 0. 4 092 147,.00 lo.,o...3 31,2 3. SO 1700 1 00.0 Non. Nil 16. S 9 1.a00 9 00 19.4 32. 3.73 1700 100.0 None Nil ZOO 94. SO129 00 3. 3 I S. 31. 3-.7 1700 2 00.0 9,000 72 0.041 Nil.41 0.41 0.45 3300 99,00 2.7.7 31 31.76 1700 200. None Nil I4Z. 000 96,4 00 3.3 6.0 32,2 3-33 1.00 2,0 None Nil - - - -- - 173.000 105.600 24.6 32.2 312 3.2 180 10.0 I12,B00 95 0. 03 Nl 9.3alost) 9.3) (et)0.1 a (o.t) 165,.000 114,.000 1. S 11. 30.6 319 * 18300 10.'0 I2',So 3 0.06 0. 02 2. 62 2. 6 2. 6 171.000 11000 12. 13. 8 31.3 3.61 1300 10.0 Non. NHI. -176.000 10,500 22,oo zz. 26. Z.4.-. -- -'- - - -* - - -. - -- - - - - - - - - - - -- - - -. -...... -............. -. 10o0 100.0 7.900 77 0.07. Nil 17..17.' 24 139,500 119,000 3.5 4.1 31.9.26 100 00.0 7o,o500 62 0.04 Nil 1,,70 1.70 1.74 130,200 106.000 3.4 11.7 2. 3.32 1300 100.0 None Nil..137.00 102,600 4.2.9 32.4 3. -79 10 200.0 6.000 67' 03 l 1. 1.0 1.. 3. 2.:2 S...73 1800 200.0 Nono Nil.....1,000 109.300 6. 3. 3R-S 1900 10 None Nil - -- - 0,00 10.800 28. ~26.0 31., 56

Table 11 Cracking Depth Data for Creep-Exposures of Rene' 41 at 1200~ and 1300'F Exposure Conditions Room Temp. Tensile Ductility Attack Depth Temp Time Stress Total Plastic Elongation Red. of Area Edge Crack Deep Crack % of Fracture Remarks -- normal thermocouple Spec, No. (*F) (hrsa) (psi) Def. (%) (%) (%) Depth (in.) Depth (in.) Surface Oxidized procedure except as noted -- R-101 1200 10.0 129.000 1.43 14.0 17.6 Nil nd -.300-inch dia, gage section except R-112 1200 10.0 132,000 1,.79 13.1 18.6.0005-.0017 nd - as noted R-142 1200 10.0 134,000 2.26 14.2 17.6.0005-.0013 nd - R-108 1200 10.0 130,000 3.67 11,.0 13.2.0014-. 0054 nd - R-128 1200 10.4 133,000 4.02 2.1 6.0.0004-.0010.023 - R-146 1200 10.0 136,000 4.70 1.3 2.0 -.036 - R-66 1200 10.0 133,000 3.49 4.7 5.3 -.022 - 0.250-inch dia, gage section R-115 1300 10.0 108,000.44 17.3 19.2.0010 nd - R-99 1300 10.0 106,000.63 17.6 17.2 - nd - R-119 1300 10.0 108,000.68 11.9 13.7.0005 nd - 1 R-102 1300 4. 5 108.000+ 1.90 15.5 14.5,0014-.0020 nd - R-95 1300 10,0 108,000 2.71 10.2 12,.9.0007-. 0021-(. 014) nd - R-130 1300 10.0 112,500 3.74 8.0 11.3.0005-. 0038-(. 035) nd - No TC on gage section R-193 1300 7.7 113,000 4.04 2.0 4.1 -.020 - 3 TC, all tied on with alumel R-189 1300 9.4 113,000 4.07 6.5 6.7 -.015 - 3 TC, all tied on with chromel R-127 1300 10.0 112,000 4.30 0.9 4.1 -.030 2.8 R-63 1300 10.0 108,397 4,.60 0.4 0.7 -.081 11.3.350-inch dia. gage section R-196 1300 8.4 113,500 4.64 6,.4 6.8 -.006 - 3 TC, all tied on with chromel R-140 1300 10.0 113,500 4.85 2,.1 3.0 -.045 - R-120 1300 100.0 86,000 5,.19 1.2 2,.9.0007-,.0042.035 2,.7.350-inch dia, gage section R-147 1300 97.3 88,000 5.21 1,.5 5,.5 -.026 - R-145 1300 10,.0 114,000 5.34 5,.9 8,.7.0013-,.0036 nd - No TC on gage section R-152 1300 9.2 114,000 5.59 2,.9 4,.1 -.021 - No TC on gage section: alumel looped around spec. R-156 1300 8, 7 114,000 5. 88 6, 2 22.6 - nd - Spec. completely wrapped; half with alumel, half with chromel, but 0,0455inches remachined before tensile test nd = not detected - = not measured

Table 12 General Surface Attack Data for Exposed Specimens of Rene' 41 Temp Time Stress Total Plastic Attack Depth Spec. No. (~F) (hrs) (psi) Def. (%) (in. ) R-15 1300 36.2 100,000 2.6 (R). 00006-. 00013 R-147 1300 101.3 88,000 5.21. 000075-. 00018 R-113 1400 8.0 85,500 6.47.00006 R-109 1400 14.0 84,000 1.36.00006 R-8 1400 14.2 80,000 2.0. 00006-.ooo00010 R-11 1400 14.0 None Nil. 00006-. 00013 R-6 1400 21.3 80,000 25 5 (R). 00013-. 00020 R-93 1400 104.0 None Nil.00020-.00033 R-13 1500 14.0 55,000 1.24.00017-. 00039 R-10 1600 14.0 None Nil. 00020-. 00033 R-92 1600 104.0 None Nil.00096-. 0011 D-3 1600 401.0 None Nil.0017-.0019 R-57 1700 99.7 13,000 27. 1 (R).0019-.0021 D-2 1700 474.0 None Nil.0034-. 0040 D-5 1700 2012.0 None Nil.0052-. 0062 R-56 1800 10.9 13, 500 32.8 (R).0007-. 0012 R-12 1800 14.0 None Nil.0010-. 0014 R-64 1800 13.6 13, 000 35.4 (R).0010-. 0013 R-91 1800 104.0 None Nil.0025-.0027 D-4 1800 1150.0 None Nil.0072-. 0078 (R) Rupture ductility 58

Table 13 X-Ray Diffraction Data from Extraction Residues of Dilatometer Aging Specimens Surface Layers of Specimen D-2 Exposed 474 Hours at 1700F interior of Dilatometer Spe --- _ ___________________________________________ _ Interior of Dilatometer Specintriormensiltometr Speimen Location: Surface Layer I Layer 2 a r3 "" ~ ~ ---------------- ---------— T Locatio: Surface Layer Layer 2 Layer 3 Standard Pattern for Indicated Phase Specimen No. and'Exposure Condition De pth D-3 0.2D Examined:.0012.0010.0011 0014 Cr0 y A103 TiN TiC NiCr 0 MC M C 1600'-401 hrs. 1700' 474 hrs 1800-5 rs (in.> _____________________ __ _____ ________ _ ___ __ _ _ _ 2 3 2 3 4 6 236 (A)'d" (A)d'I I (A)" (A)'d' d" I ( -L A d (A)'d' t (B)"d-d - () 3.628 20 3.63 74 3.479 74 3.628 20 ------------------- T T3 —74 3 ^ -— ^ - — d — I — "^ I — -- ^ --- I- — "d" I "d" ( (A)"-d" I (B) -d"" I (B))""d" 3.610 20 3.616 30 3.610 10 3.575 20 3.268 to10 3.2:,37 5, 3.12 5 3.237 5 3.20 3.105 3.05 5 2.9926 5 2.93 30 2.796 10 Z. 913 10 2. 77 in 2. 772 to 75 5 2.666 2.660 30 2.67 100 2.68 w 2.693 5 2.650 80 2.649 40 2.554 20 2.552 92 2.54 m 2.546 Z.6 82.486 40ms 2.48 96 249 s 2.50 100.48 2.48 30.45349 ~~~2.2 467 8040.45 2.46 vw2.535 5 2.32450 9230 2.448 90 2.399 s 2.395 40 2.406 70 2.394 15 2. 382 1 0 2. 341 40 Z.379 42 2.262.261 20 2.63 2.53 10 2264 40 2252 20 Z.256 10 2.26 12 2.235 5 2.13819 50 2.18 5. 2.189 40 2.196 60 2.190 20 2.73 5 2.,17 38 2.165 <1 2.406 5 2.166 20 2.166 10 2.16. 2.135 u 2. 130 40 2.135 80 2. 126 80 1 1 2.120 100 2.114 100 2.1 40 2 12 2. 105 2.085 10 2.085 100 2.07 35 U-I0 2.020 10 2.026 5~~~~~~~~~~~~~~~ 2.064 100 2. 05 9 2.62 0 2.068 100 2.069 100 2.065 80 1.960 0 1.959 n 1.953 10 1.961 20 95 30 1.894 20 1.875 w 1.892m 1.892 40 1. 901 0 1 L.810 40 1.817 10 1.811 20 1.818 30 1.82 39 1. 850 w 1. 810 s L. 846 5 1841 0 1. 803 5 l 740 43 1.753 vw 1. 785 mw 1.789 10 1.819 20 1810 20 1.738 10 17421. 70 5 1.791 *788 1.6685 5 1. 69 w 1. 692 vw 1.692 10 1.714 10.694 10 1,6662~~~~~~. 670 30. 67 90 l. 67 w 1.667 5.665 10 1.~666 100~ 1. 669 20 1.665 70 1 632 vw 1.642 5.634 5 *1.616 10 1.. 60 8 Ls 13.,*615 r-w 1. 616 20 1.622 5 6 1 1.600 20 1.602 10 1 5 13 1. 6 01 8 164 1 0 6 I~ ~ ~ ~ ~ ~ ~ ~ ~~~~~~~~.4 1.608.6021.609L0 1.52Z3 20 1. 510 7 1. 528 M 1. 552 row 1. 545 1. 553 10 1. 497 80 1.497 60 1. 500 20 1. 50 M 1. 528 w 1. 500 VW 1. 522 20 1. 525 10 1. 478 20 1. 47 80 1.481 1. 486 vw146 5 1.463 30 1.464 10 1.465 10 1.46 12 1.443 1.431 1.447 10 1.446 10 1.430 40 1.432 10 1.430 30 1.434 10 1.43 40 L. 385 1. 374 10 132 10 1.40 10 1.404 3 1.34 10 1.0 344 152 5 1.375 10 1.374 480 - - - _ __ - -- _1._-_ _ _ _ 30 20 1.31 10.9 2.9 3 Cra03 Cr03 Cr03 TiC? 1.28 M M6C MC MC TIN TiN TiN MC M6 Ni~~~~~~r 0? Cr 0? ~~~~~~~~~~~~~~~~~~~~~23 6 M23C6 23 6 NiCr2 4 2 3 TiC slight TiC 2~~~~~~~~~~.84 23 V- AO1203 C? 23 6 "d" interplanar spacing I - relative intensity (A) - 57. 3 mm camera (B) - 114.6 mm camera t ---— I —-I --— I____ II___ _ I__ ____I_____ II _ __ I__ _,______

NOTE: 0.025 INCHES MACHINED OFF GAGE SECTION DIAMETER AFTER CREEP - EXPOSURE 4 o \ ~O ~ i-13 NG 0 \4.0 ),[" (/ 0 0\ o0 -ff ili 7 -, --. _i' I~~5 i APPRPOX. 1I 5 a r - 2 -t 4.0 TENSILE AND CREEP SPECIMEN GRIND BOTH ENDS FLAT AND PARALLEL 4.0 0 D LATOM ET ER SPECIMEN o o 0.010 R- <1 —.08 —- 8 0 < —. 2.16 ----— > * l97 TYPE-W IMPACT SPECIMEN Note: Center of the Compression and Impact Specimen to Coincide with the Center of the Creep Specimen. Figure 1 Details of Test Specimens 60

Dial Gage Brass C onnec tor ---- Brass Cap 28 gB^ ^ ^J ^G —-Slit To Potentiometer Brass Con.-/_,,X/Brass Connector / / /Positioning 28 ga. Lip //1 / / To Potentiometer // / / — Outer Quartz Tube H o loHollow Quartz Tube Push Rod Ceramic Spaghetti Slits Quartz Connector Measuring Thermocouple - 18 Gage Specimen Not to Scale Tube and Gage Support System Not Indicated Figure 2 Method of Attaching Thermocouple For Long-Time Dilatometric Aging Studies 61

200 Code 190 &j-v. 14000 Code 180 -V-1400 C) __________ 0-100Exposure 0 — 0 —1600 170 — A —1700 Temperature P _01800 _ ___________-0 —1800 1o 160 11800~ e g = = _~~ L Y < I 00Oa~16000 0 150 0 50010001500200018000 1 140. — 4E Time17000. 3 130 -o _-_ _ PA U) 120 _________________ 0'4 1 ____________ __________' T140 1400 "R" Condition ^ 130' —- 1950~- 1/2hr+AC - 5) -r 1400~-16 hr+AC H~ ~4 120 Q) 0\ 14 O __ 9-o 6 600 N 1800 _________ % =. A 1 000 0 - vS i 90 - -016000 30 (1C 1, 20 + 800016000 o 10 -' 1700' L~ ~ ZO ^ —----------------------------- ----- - -- o t10 IAi A8 I 0 500 1000 1500 2000 Exposure Time - hours Figure 3 Effect of Exposure Time in Unstressed Exposures on Room Temperature Tensile Properties of Rene' 41 Alloy. ("R" Condition) 62

.0050 - i j1800~F i.:\800'.0030 ~,0025 ___ ^ ~ /.^D-5 1700~F.0020 ------- --- - / J/ i-3 J*5'F Thermal Expansion Effect.0015 - yiF —---- -----.0010-~ -— _________L_________.0005 0 200 400 600 800 1000 1200 1400 1600 1800 2000 Time - hours Figure 4 Shrinkage versus Aging Time for Dilatometer Exposures of Rene' 41

+.005 -- ] -- - -- 7 - I -- + 005 1IF /24,000 psi 1600~F L / / 1700~F +.004 1 — --- +_004 _~ -- -- --- 12,000 / O psi ^ +.003 -- - -- -- - +.003 +.002 —'43 — -- --- -- --. 002 +.001 +.-001o / —--- o0 Unstressed Exposures'WI ___ --- ---- I'r_ Ustressed Exposures C: In Conventional Creep Unit i 0 2 I::::=::::''"' -- -'^ —-^^^^'"'T^1" -- In Conventional Creep Unit In Quartz Tube Dilatometer -- _ _ -.001 --- -' -- -- -- - -.001 In Quartz Tube Dilatometer -.002 -.002 0 20 40 60 80 100 120 0 20 40 60 80 100 120 140 160 180 200 Time - hours Time - hours +.005 -\ - ^~~~~~~P~~~~~~~~~ ~7,7500 psi 1800OF +.004 -- - — 46,000 psi U | +.003 —-- - I +.002 ~ +.001 ^P I I~~1 QUnstressed Exposures In Conventional Creep Unit -, 001 In Quartz Tube Dilatometer - 002 0 20 40 60 80 100 120 140 160 Time - hours Figure 5 Comparison of Shrinkage Observed in Dilatometer Aging With Data Taken in Creep Units

.010......... 12000 F.009.008 /eat Treatment /950~F- l/2 hr +AC'. l/ 1400~F-16 hr +AC.007 I |114, 000 psi.~.006'.006 ~~ r l | I Heat Treatment,.005 2150~F - 1 hr + OQ bo o i/ 1650~F - 4 hr + AC 0.004 75, 000 psi.003.002 0 20 40 60 80 100 Time - hours Figure 6 Negative Creep Observed in Rene' 41 at 1200~F During Initial Survey Tests (Ref. 1) 65

40 --- --- 0 0 30 - - -- -___ ______ ____ *H 4o 0 F 7 c 0 u U-10 hr g ^ 20 U ^ u0\0 ICa 0I U0 g 100 Specimen Size: 401 0. 197x0. 197x2. 16 inches 0 200 400 600 800 1000 1200 1400 1600 1800 2000 Exposure Temperature - ~F Figure 7 Effect of 10 Hours Unstressed Exposure on Smooth Bar Charpy Impact Strength of Rene' 41 0 -Ul —---— timate Strength ------------— T 100- T 10 Hour Exposures ) A 0 Smooth Charpy Strength X < (Average Values) " 401 Hour c^ <D 0^ I ) I I Exposure 0 200 400 600 800 1000 1200 1400 1600 1800 Exposure Temperature - ~F Figure 8 Relative Efficiencies of Tensile and Smooth Bar Impact Strengths in Showing Damage to Rene' 41. 66

.20 ---- 1400~F' —- - --------- __ __ __ __ __ _ _ - i i"83% of Rupture Life ----- Creep-Exposure Followed by Indicated Re-Heat Treatment 15 U c R- 72.05 - - _- I - - - R- --- -61,000 psi R-150 1 84, 700 psi 0 10 20 30 40 50 60 70 80 90 100 " ---------------— 6, 0 -- — Time - hours 50 -47s250 Ultimate Tensile 50 250o. ---- 0.2% Offset Yield All Specimens Remachined S 40 2 n -Elongatio Before Tensile Test 40 200 Elongation OI I Reduction of Area RHTR: 1950~F - 1/2 hr+AC l30 h 150 1400~F - 16 hr+AC - | E20 os 100 0 & 10 s 50 c (C 0 0 As HTRAs Exp. R-150 As Exp. R-172 to 3. 5% RHTR Approximate Level of to 20% RHTR Creep in 10 hrs, Properties After Creep- Creep in Exposure 100 hrs. Figure 9 Effect of Re-Heat Treatment on Tensile Properties of Rene' 41 After Creep-Exposure at 14001F

R-171 5 24, 000 psi.20 600F --------- - 34% of Rupture Life /- 1700F Creep-Exposure Followed by / ~~~~~~~~~~~~~~~~~~~~~Creep-Exposure Followed by Indicated Re-Heat Treatment Indicated Re-^ ^ ---— ~ ~ ~ ~ ~ ~ ~ ~~~~~~~~~~~~~~~~~ndcae / S e e-Heat.~; / I I I I F I I I I I I Treatment 15 j l0 --- -- -— I- -__-`-_ _ —__ - - - - 10 - 0 1-4 R-~~~~~~~~~~~~~~~~~~~~~~R185 0 3 R-448 39,000 psi 23.500 psi 0 -74% of Rupture Life 74 of Rupture Life 05 ------ __ -----— _ —----------------------------------— ___ —---, -- ---- -- ^ - ---- ---- o R5-151 23,000 p si.05 i 05 *./^ 05 -'67% of Rupture / 00.^ >^ ~~~~~~~~~~~~~~~~~~0 Life 00 0LI"II-~-' ^ I I --— I —------- I ----- I ---------- O ^ - -------- 0 10 20 30 40 50 60 70 80 90 100 0 10 Time - hours 50 250 ----— Ultimate Tensile All Specimens Remachined 50 p,250 ^ o I ----- 0.2% Offset Yield Before Tensile Test d 40 200 — Elongation RHTR: 1950~F - 1/2 hr+AC <S *s'II i-Reduction of Area 1400~F - 16 hr+AC o 30 I 150 w 20 2 100 -u*. a) 00 i0 50 So 0 0rl 0 AsHTR A AExp. R148 As Exp. R-185 R-171 As Exp. to R-151 to 6%o RHTR to 7.5% RHTR RHTR 3.6%Creep RHTR Creep R Approximate Level of Creep in in 10 hrs in 10 hrs. Properties After Creep- 100 hrs. inAhs LExposure A Figure 10 Effect of Re-Heat Treatment on Tensile Properties of Rene' 41 After Creep-Exposure at 1600' or 1700~F

i 74% of Rupture Life 25 —--- 1800~F — / Creep-Exposure Followed by / Indicated Re-Heat Treatment.20 1 Time I hor ts 5 o I I ^ As HT sEpR-5sEp R-197 7300 psi 0 ir / o O "69% of Rupture Life R-155 12600 psi C0' _.05 0 -'85% of Rupture Life- {-' __' —----. —.. \O^~~~~~~~' / >^^^^ ~~~~~~~~~~~~~~~~~~~R-124 7000 psi I/ -^^^^~-' -57% of RuptureoLife R2-16 7500psi -74% of Rupture Life 0 10 20 30 40 50 60 70 80 90 100'3 Time - hours es 50 ~ 250 0 2 All Specimens Remachined R Condition R2 Condition 40. 200 Before Tensile Test 40.c',00 ~ —---— Ultimate Tensile' f -[ {[ —-- 0, 2% Offset Yield RHTR: 1950'F - 1/2 hr+AC 30 150 Elongation 1400'F - 16 hr+AC C II < I 1I —Reduction-of Area RHTA2:2150~F - 2 hr+AC'- 20 ^ 100 1650~F - 4 hr+AC 00 i 50 [jjR2: 2050~F - 1/2 hr+AC c - ^ ^ 1650~F - 4 hr+AC o o'l o __________________..i__..__..___ AsHTR As Exp. R-155 As Exp. R-197 R-124 R-116 As HT As Exp. RZ-16 to 5A% RHT toRHT RHTA RZto 2 RHT0 to5 HR Approximate Level of CRHTe RHTA RHA A2 Creep Properties After Creep- 100 hrs. 100 hrs in 10 100 hrs. 100 hrs hrs. Exposure Figure 11 Effect of Re-Heat Treatment on Tensile Properties of Rene' 41 After Creep-Exposure at 1800'F

a 10 hours Prior Creep-Exposure 100 hours Prior Creep-Exposure | 1.1__ _ __ I —-- _______Code s u 1 2 A\ j -- --- Original Properties - > | ~ 1.0'r -—' - -Condition "R"' c. - _ ---- Unstressed Exposure *'8 _ __ -_ - - -- Creep Exposure -Heat.' |-,''' vI..|Treatment "R". Creep-Exposure + Re-Heat v Treatment "A" o 1. 7' _' All Specimens Re-Machined Before Tensile l. 6 - - - _ — 0_.... _- -...............Tests l.c 5 - — \ --- Heat Treat Code H^ ~ 0^-S *\"R": 1950~F - 1/2 hr + AC 1! -- - -----. 1400~F - 16 hr + AC o i u) L. 3- -\ -- ----- ----- "A": 2150.~F - 2 hr + AC - ~ *... \ 1975~F - 1 hr + WQ a. 1.- -- 1400~F - 16 hr + AC h 0 1. *! —— ": — -- I d 8 A 14 | * —-- O. O:3-. - 1. - E 1.0 -* 6 0)I01 I I 1.3' - Unstessd e 10.1 o. __ 0.1 3 C-Unstressed exp. Tm a Creep Exp. + RHTof.,-<1_ 41 An S s u Ait o Origina l i 010. _rigina ___....7 H0 0 N'.... 0 __J l/ ff Cr e p E xp.A?,-, [ -0.......repEx. 3 1 x ___ OUntese R' —Ip],n -sree O ____ Creep Exp. =:.4 - _ _ _ -_ _*2 Prior Creep:: - Prior Creep: —t[ —O,Z 24 Z L8;4 %f ~1 3.5% 6.Z% 3.6% 5.3% — - 19.8% 7.450 8.5% -- 0 -- 1 0 - J ---- — 41 0 II J 400 1600 1800 2000 0 1600 1800 000 Creep-Exposure Temperature - F Figure 12 Effect of Creep-Exposure on Room Temperature Tensile Properties of Rene' 41 And Subsequent Ability to Recover Original Properties by Re-Heat Treating and Resnachining 70

90 1400F I --— \0'F Estimated Original Rupture Life - o 0+x,,, A (8. 0) NN| " \(3.7)-,,"% 80 - -s \ (25)-\_ Expected N. a x ilst Exposure Life for 2nd Actual Life Exposure o Exposure 2nd Exposure 70 o N'_-_ —-" S C | \\' <^Y \ Total Life N (18. 3)0.. " (32.8) 60 (18. ^ (14.5)Z 8) 50 1600~F Est. Original Life 25) 40 -2 —-- --- 0.' 7' 9)^a t302) - -—.) u) /( 6 P>\>\di(NN37.1) 20 - - Expected Life ( ------ \ - Total2nd Exp. kL Actual Life 1st Exp. Life 2nd Exposure 10 ____L_________________________ _____I________________I________I____ 30 30 1700~F (6.3) -< (28. 4)'-x.. "~m~...~"a..(34.7)'~,~_______~ | ~ -'~~ Total Life 20 ----- Expected Life/- / Actual Life o Exp. (27) 10 20 1800*F + 1st exposure O 2nd exposure A total life x expected life 2nd exp. 4 Est. Original Life__ ( ) Elongation 10 (34) J 1 Heat Treatment "1R" 0 20-, ~1st Exp.!^ ^ 1950'F-1 /2 hr+AC 18 |..._ _.1) |14006F-16 hr+ep AC_____' 10 ^ —_^ ^'!^*^^^^. ^_ _Total Life o Expected Life. (8.4) 5 Et 2nd Exp. Actual Life... 2nd Exp. 1 10 100 Time - hours Figure 13 Effect of Re-Heat Treatment "R" on Rupture Life After Initial Creep-Exposure 71

5. 0 Initial Exposure and 0 Re-Heat Treatment Both in Condition "R": 1950~- 1/2hr+AC 0 1400'- 16 hr+AC 0 4.0 -hr- - - - - 4,70:ex......... Ist exp. A. In w"~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~r L 100 hr m^~~~~~~~~~~ \(d ~~~~~~~~~~~~~ \Istexp./ k 3.0 - - 15, h 100 hr \0 -.... Istexp.\ coverabae ~1H 1 ( S If all life were 2- o 1 0 Ist - " 0 1 hr 1st exp. 0 /.. A. < c o i f^ ^.^ I'^S^- ^^^^?A'? ^ Unaffeted tlc~l o a, recoverable x o'0 1. -10 hr- 1. — e. _ 0~Lst exp0 cd E la 0 01. 0 0 -0 f Life F raction [ 100h hr / I o o Aule Obeyed 0 0 0 R O y Ist ex. o Actual Exposreep TimeRe 0 -- — Lif. | L- -----— e -Freact 0 0 0 ule Obeyed Ist exp. ef Life Fraction Ee 0 1400 1600 1800 0 1400 1600 1800 0 1400 1600 1800 0 1400 1600 1800 0 1400 1600 1800 a) Exposure Temp - F b) Exposure Temp - ~F c) Exposure Temp - ~F d) Exposure Temp - ~F e) Exposure Temp - ~F Figure 14 Effect of Re-Heat Treatment to the "R" Condition and Re-Machining on Ratios of Actual Behavior to Theoretical Behavior of Rene' 41 After a First CreepExposure of 10 or 100 Hours at 1400', 1600', or 1800'F.

.40. - - -------- ------ - --- ------ ---— R-157 1800-F 12.600psi R-154 1700-F 23.000psi yi r I' IC Re-Machined: 0,025-in, from Dia, I I~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~I, - 1st Exposure.' f Then Re-Heat Treated: 195-1/2 hr+AC 1600F 39.000 psi.30__ and1400 -16 hr+AC - - — __ I_ __ _ __ and Re-Machined: 0.025-in. from Dia, -— X 2nd Exposure to Fracture r / Stress and Temperature Identical I / For First and Second Exposures I I.zo i I / I I II I ^J~~~~ ~ ~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~ I / I I I I 0 / I I --- -- ---- — /-I- - -I. —- --- -.20 I I I r I I / I I /r LA I I 103 I / I, I / /l I. / /i/ / / -I- / -' —- - I 153 400F 84.700'ps / / /,- -- 0 I 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 2 9 3 /'9 II/ / /, -- - *10 - - - - -'% - - - 51 R-153~ 1400F 84.700 psi 0 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 Time - hours Figure 15 Effect of Re-Heat Treatment and Re-Machining On Creep of Rene'41 After Initial Creep-Exposure For 6, 9-10 Hours

.40 l R-178 1600~F 23,500 psi 1400-16 hr +AC and Re-Machined: 0. 025-in. from Dia. 1 / I -— kZnd Exposure to Fracture i j 1 rR-179 1400~F 61, 000 psi Stress and Temperature Identical / / 1 I I, For First and Second Exposures / - - - 1 3i I3<';/' / -'~~~~~~$ i I - i... I, I I // - hours Figmd1Expos ure 16 EFrfctue/ o R -170 1600 F 23, 5000 psi "^' r/';S "R3" Re-Heat Treatmentil V /I l \ 19506-1/2hr+AC *~ /~~~~~~~I/ II,I 16008 1/3hr+AC r.0 1 I - - - I /...... -~-. -. - /- _ _ _ _/_ __....... t i H= i I -/! - 1 I! i ) I I i ( i I I ) I I I / I,, I'10 - -!-' ~ I -t —----— ~ —— j, - 11~" -~ _______ 0 10 20 30 40 50 60 70 80 90 100 110 120 130 140 150 160 170 180 190 200 210 220 230 240 250 260 270 280 290 300 Time - hours Figure 16 Effect of Re-Heat Treatment and Re-Machining on Creep of Rene' 41 After Initial Creep-Exposure for 100 Hours

14 (33) 13 ((35) (32) -1800~F A'O (20) 12 __________ (N 11 1o Code —------- (34) \ (19) 10 Cd (9 0 ) A(21) 9 _ ( ) Approx. Elongation __ (9) \',^ (19) 9 O "R" - Original Life (As Treated) (9 \. --.^ | |+ "R" - First Exposure \ F8 0 "R" - First Exposure + RHTA \.. (41) _____, o A "R" - As Treated + RHTA " (23) Zl21) o;B - 7 V "A"- As Treated. l). o Heat Treatment Codes ^ ^ ______R' F.. "R" First/ is "R" First Exp.......... RHT. Exposure + RHTA; RHTRA Re-Heat Treated RHTA only; RHTA- __ R' As- A 5 T "A" only "R": 1950~ - 1/2 hr + AC Treated 1400~ - 16 hr + AC "A": 2150~ - 2 hr + AC 1975' 1 hr + WQ 3 - 1400 - 16hr + AC | 0 1 100.0. 0 10 100 1000 Rupture Time - hours Figure 17 Effect of Various Re-Heat Treatments on Rupture Life of Rene' 41 at 1800~F

40. ---— 01vst Exposure Heat Treatment Codes Then Re-Heat Treated Initial Exposure in Indicated Condition and ReMahined in fo D RHT- Re-Heat Treated to Indicated Condition and Re-Machined: 0. 025-in. from Dia. 1950~-l/2hr+AC - — X2nd Exposure to Fracture 1 0-1 hr+A C Stress and Temperature Identical - hr+AC For First And Second Exposures - A: 2150'-Zhr+AC _____ __________ 1975 -lhr+WQ 1400 -l6hr+AC RR-166 12,500 psi I R+RHTA R-157 12, 600 psi R -?6 I, 00 siR - 13Z 93500 s * lZ,*600P511 R-126 12,500psi I,<R-l, |,5 0 psi | ( j AU Creep-Exposures at 1800~ F / _______~~~~~ I _(? ___;RT ____________________~_____. __ _ _i _ __________ ___ i 1.5 / I I cG IRHTA / 2,I II I 9 J> I: /|1 / R-7 3~~ ~ ~~ I -- ___________________________ / 1 _ _ ___ / - ___ /9.500 psi 0~~~~~~~~~~~~~~~~~~~~.20i / - I, f ^^-^ ^ yI R F / ________ -— ___ _ __ _ -- /1 -- - - - - --- I " - 1 0 10 20 30 40 50 60 70 80 90 100 110 120 130 140 150 Time - hours Figure 18 Effect of Type of Re-Heat Treatment (Followed By Re-Machining) On Creep of Rene' 41 After Initial Creep-Exposure at 1800'FTAFor 6.8-33 Hours.10.011 Ie /~~~~~~~~~~~~~~~~~~~~~~~~~~~~~ 7 09~~~~~~~~~~~~~~~~~~~~~~~~~~~0 -1.001 0 10 20 30 40 50 60 70 80 90 100 Ito 120 130 4 5 Time -hours Figure 18 Effect of Type of Re-Heat Treatment (Followed By Re-Machining) On Creep of Rene' 41 After Initial Creep-Exposure at 1800"'F or 6. 8-33 Hours

.40 --- 1st Exposure Then Re-Heat Treated and Re-Machined: 0. 025-in. from Dia, -— X 2nd Exposure to Fracture Stress and Temperature Identical For First and Second Exposures Heat Treatment Codes Initial Exposure in Indicated Condition RHT-Re-Heat Treated to Indicated Condition __ _____________ ~R: 1950~-l/2hr.+AC___ ____________ _ _______ *.~~30C-~~ 1400~-16 hr+AC I A: 2150 -Zhr+AC I^ ~1975'-lhr+WQ I 1400"-l6hr+AC I I R-121 All Creep-Exposures at 1800'F R-14 I U" I.5 / /o R-168 7,300 psi 5 I I I 4J / I c - i /RHTR R -j / / /?R-133 J 8,000 ps l 7,300psi 0 /R- I 0'7,500 psi 00 8,~~~~~~~~000 psi d~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~ O~~~~~~~~~~~~~~~~~~~~~ 0 / / HT o~~ ~ ~~~~~~~~ / / RHT — i" Initial Properties / A 7,300 / / psii ~~~~~~ps7____________ ______ "__________________ /_____ ^-_____________________________ // T BH^ ^ —-~~~ ~~ —-— 7 0 100 200 300 400 500 600 Time- hours Figure 19 Effect of Type of Re-Heat Treatment (Followed By Re-Machining) On Creep of Rene' 41 After Initial Creep-Exposure at 1800~F For 100 Hours

i ~ ~~~~~~~~~ 1 0. 30. 1600iF - I, I~~~~~~~ ~ ~ i'' I I I i;' j!~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~'..R / f,,/,,,,Re-Heat Treated I Only a ~~~ ~~~ ~~~ ~~~ ~~~ ~~~ ~~~ ~~~ ~~~ ~~~ ~~~ ~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~ 0.3 20 I600OF' t i....;''' — 3, pi _ j-18 14 I __ I _: R H T R+ 1 - - -,/t _ _ _ 00 r 1 I I~~~~~~~~~~~~~~R- 4 /, 39,000~~ ~~~ ~~~ ~~~ ~~~ ~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~ 14001-16 hr+AC ~ ~~~~~~~~~~~~~~~~~~ R- 187A ~ ~ Only i/ and~ Re-Machined - i I I __ 0 1 2 3 4 5 6 7 8 9 10 i 12 13 14 15 16 17 18 19 20 21 22 23 24 5 26 27 Time - hours Figure 20 Effect of Re-Heat Treatment and Re-Machining on Creep of Rene' 41 After Initial Creep-Exposure at 1600'F I'!,~~~ Ir~~~~~~~~~~~~~~~~~ 0~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~, 10'~ ~~~~~~~~~~~~~/j 39,000~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~ ~~15"12h+COl s,, 1 1400"-16~~~~~~~~~~~~~~~~~~.!+A I t R-t83A 39,00i?s f -- I ~~~~~~~~~~~~~~~~~~~~~Only Time ~~~ ~~- hor Iiur J0 1sft Exosue-eareamn an-R-"ahnn Znd Cre fRn'4AtrIiilCeExposure Jt!.~'.

* ~ ~~~~~~~~~~~~~~~~~~~~~~~~~~~~~ "A ** ts^ ^ ^"'^ A -— ^ tv T o~~~~~~~~~~~~~~~~~~~~~t ^ ^ - v. >^^ ~*t'^^*<>.<~^^">^^***WY 4,~~~~~~~~~4 I. 41. 1~t* *;Ho?%?s;>~... —— \ y } f~~~~~y~f-^yfif9 ^;^S^:l ^ ^V^'^^*^^ *.'^.^<^^1<^~~~~~~~~ x4 ^><'! * a.*- ^ ^^p^ ^ tA Figure 21 10OOx Figure 22 10OOx As Heat Treated "RI' Condition Specimen R.-104 Exposed "Without Stress 100 Hours at 1800~F v^^7 ^' 1 | l * * * t t.' *~~~~~~~~~~~~Iv Fiur 23 I 0O Jiur 24 I i>

4 Figure 25 1000x Figure 26 2200x Optical Micrograph of Spec. R-150 Electron MiGrograph of Spec. R-150 Specimen R-150 Creep-Exposure 10 Hours at 1400~F to 3..52 Percent Deformation, Then Re-Heat Treated to Condition "'R" ~'i:..~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~. it it'~~.......... ~ ~ ~ ~ ~ ~ ~, Figure 27 2200x Figure 28 2200x Electron Micrograph of Spec. R-172 Electron Micrograph of Spec. R-148 Creep-Exposure 100 Hours at Creep-Exposure 10 Hours at 1600~F [400oF to 19.0 Percent D:eformation, to 6.20 Percent Deformation, Then Then 1Re-H~eat Treated to Condn. "R" tie-Heat Treated to Condn. "IR" 8O

Figure 29 2x Figure 30 2200x Electron Micrograph of Spec. R-185 Electron Micrograph of Spec. R-151 Creep-Exposure 100 Hours at 16006F Creep-Exposure 10 Hours at 1700~F to 7, 38 Percent Deformation, Then to 3. 57 Percent Deformation, Then Re-Heat Treated to Condn. "R" Re-Heat Treated to Condn. "R" ~. Figure 31 2200x Electron Micrograph of Spec. R-155 Creep-Exposure 10 Hours at 1800~F to 5. 34 Percent Deformation, Then Re-Heat Treated to Condn. "R" 81

Optical Micrograph of Spec. R-197 Electron Micrograph of Spec. R-197 Specimen R-197 Creep-Exposure 100 Hours at 1800~F to 8.06 Percent Deformation, Then Re-Heat Treated to Condition "R" " ~i t /,t -', f:^.)':"'rt 1^t},1 > f /:'. *1~~'A A-~.:. * Figure 34 1000x Figure 35 2200x Optical Micrograph of Spec. R-124 Electron Micrograph of Spec. R-124 Specimen R-124 Creep-Exposure 100 Hours at 1800~F to 3.46 Percent Deformation, Then Re-Heat Treated to Condition "A" 82'~~6~;1: ~~' 9 ~~~~~s t I~~~ irl C -f,~ Figure 34 lOY iue3 Z Opt~icf1 al Micrgaho pc -2 lcro irgaho pc -2 Specime R-2 Cre-xpsr 10 or t100 o34 ecn Deomton hnR-Ha ratdt odiintf "4 o.LP i~~~~~~8

5000 ^o-D 1800~F //^./i o 7 / O r 3000 ~ O O/ 0 1600~F 2u000 - - As Treated "RI P 10I 0 As Treated_-R211 <a)0 0 0 1600) Exposure A 17000 Temperature i 8 O O 1800~ OF 0 L I I 0 200 400 600 800 1000 1200 1400 1600 1800 2000 Time - hours Figure 36 Effect of Aging Time and Temperature on Average Gamma Prime Particle Size in Rene' 41 83

50 Code C*rd As Treated "R" yX C As Treated "R2" *.40 -- "R" "R2" d ^0~~~~d | |o a 1600~ Exposure E~~g rt'~~ _ I a 1700~ Temperature S. 30 F Z | n o 1800~0 F 0.30 oo o --- i; co 2t ~ - 1700oF.,20 - _,.20. I....... --—. —. -.... —-- aI*........ _ - _1800~F >. 10 0 500 1000 1500 2000 Exposure Time - Hours Figure 37 Effect of Exposure Time on Volume Fraction of Gamma Prime in Matrix of Rene' 41

140 130 - z7)\0 ('37) (.27) 120 \ III \ (. 26) 21________ \(.2l) 1 10 (.23) 0 * A(.25) \0 (.22),100 ) (.19) \ (.24) ~ (,21) 100 o \(.14) _A",~~ A (.21) A (.19) ___0 _ _ _ __ (.18) 90 80 ____ Pb>4~~~~ oCode ~ I O As Treated "R":i X As Treated "R2" O "R" "R. 2'" e.7 0 0 0 1600 ~ Exposure o A 1700' Temperature O 1B 1800'~ ~F ( ) Gamma Prime Volume Fraction 60.________- L - J —---------------- — _........,,.............. 60 50' 0 1000 2000 3000 4000 5000 Average Gamrma Prime Particle Size - A Figure 38 Effect of Gamma Prime Particle Size on Yield Strength of Rene' 41 85

140 r......... 1.......... / /.....Range for 1o2 - — / -------- - / ----- ------------ --- "RAs-Treated /120 /"R" Condition * | /d As Treated "R" /As Treated "R2" 50 i / /A / ~ ~~~~~0 L16000) "R", After 14-2012 Hours 100 A 17000 Exposure At Figure39 Co ti on / / / F tion O 18000 Indicated Temperature If t 16000) "R2", After 14-104 Hours of. Rene' 41 After A Unstese Exposure At oltn U 8 1800~ Indicated Temperature 90 // f, Volume fraction of freely dispersed -4?~~ / / / ~~~~~Gamma Prime >-1 80 70 Mill Annealed Condition 60 (Solution Treated) __._____..______ 50 0 1 2 3 4 5 6 7 8 9 Meiklejohn-Skoda Volume Fraction Parameter - 0.82 - f'/ Figure 39 Correlation of Meiklejohn-Skoda Volume Fraction Parameter With Room Temperature Yield Strength of Rene' 41 After Unstressed Exposure 86

...~,~,.......;:~~~~~~ ~ ~ ~ ~~~:,,:: i: ii.'.~...::: I:,::::~,~'::lil:-:i:;i ~bj- ~ -Fi g u r e 4 0 8 200x Figure 41..'..;::.........:Il::i:;l:::::: l':'_':'i~~ailiii~ili I:' " i::?~:~i'::;;- ~'::: < ~ i:;~:':!~ 1 XI bill:-il:,:::::i'::!l:':~'i-i~~~~~~~~~~~~~~~~~~~~~~~~~~~~~i=~:;il~~~~~~~i~~~i:~ii?....:......'~:'...:;~:/::........ ~l': I:::i "'':,::-,"':e'-i:....': "'::!'" ~;:i " ~:'~'':':!''':;':'?''' ~~~::~~~~i:-i::~:;:::::i~~~~~~~~~~~~~~~~~~~~~il:::::::i;::.~~~~~~~~~~~~~~~~~~~::;i:~~~~~~~~~:::.::: ~~~~~~~~~~~~~~ii,:~~~~~~~l,:;i::::: ~~~~~ ~ i~~i!!: Figure 40 4800x 35gur 41400x Coldo elc fSe.R-2Carbon Extraction Replica of Spec. R,-32 Carbon xtractin Replia of Spc. R-32 Collodion Replica ofSpec. R-104:::ii,:::;i;::::~8

_ZMl:~~. -j - <; t: ntv W /*.9* _ Figure 44 8200x Figure 45 200x Collodion Replica of Spec. R.-104 Carbon Extraction Replica of Spec. R-104 -U-~~grl( *:~; A Figure 46 22 ZOOx Figure 47 220;X CarbonExtraction Replica of Spec. R-104 Carbon Extraction Replicaof Spec.R-104 Spec. R-104 Exposed Without Stress 100 Hours at 1800.. 3 b > # We; I~: s:..,;:., rn- 0 >';*3 > 0' 1.t:0+ { e 0;} *zt':.. outlined area) Spec. R-10 4 Exposed Without Stress 100 Hours at 1800~F

Figur'e 0' o;48 4800x Figure 49'. Carbon Extraction Replica of Spec. R-104 Electron Diffraction Photo of Selected (Diffraction obtained from Area of Figure 47 outlined area) i,...,..*.. - f^ ^.^-.*W 1' - wN.......... Figure 508 46 00x Figure 51 1649'00 Electron Micrograph of As-Treated Electron Micrograph of Spec. R-124 Condition "R" (1950~ F Solution) Re -Heat Treated to Condition "A" o(2150F Solution) After 100 Hours Creep-Exposure at 1800~F 89 89

10 Hours Exposure 100 Hour Exposure 210 ——. 190 -- o a 170 —---------- \ \ — 180 \ 0 ^ 170 -------- 160' - I) 0 150 \ 130I \ 140 I I. 0 - I I ^ --- I - -- - - - - J-% I 130 -- -. I l0 -- CODE 100 --- Remrachined (Ref. 1) ( 100 - V v — 0 — Not Remachined H g 90 _ l As Treated o 90 30 g 20 -, — — 0 —~ — ^~> - -^ —-— P ~^^ — - 0, 2 _ — o — __ ____ o~.o S I o — 4-_i 10 o, L 3o 0 200 400 600 800 1000 1200 1400 1600 1800 2000 0 200 400 600 800 1000 1200 1400 1600 1800 2000 Exposure Temperature - F Exposure Temperatuie - ~F Figure 52 Effect of Remachiiing on Curves of Room Temperature Tensile Properties Versus Exposure Temperature of Rene' 41 Exposed Without Stress for 10 or 100 Hours. 90

10 or 100 Hour Unstressed Exposure 10 Hour Creep-Exposure 20 2 0 s_ __10 — + -10 -' S I:' / I ~\' _ _ _ _A _ 3 2 3ot RemacAined Re-Machined +0 o + — * 1% Creep C a E~-~ Code 0o d Reachihined R A d 1 ~o | ||30 3- 0 —A —4 — 0 - — | —— A —A Not Re-Machined: ~.... - ~0E-po — u — Unstressed 100 hr Exposure 10i i..m 10 — O —- 1-% Creep O Diffren' sp d Remac hined, 0 -— A —- 4% Creep ^~~ e2~~~-o 1 a0 - 10A —--------- 0 0 - t~- ) 6~ _' I j I - i k a 20 20 20 - 20 a,^~~~~ II 1 0 hours 4 +10 10 I 0' +' — At Unstressed Q, c —A i-,~~ o (S 20 (- — 2. —— CreepI4......A_ 4 % 4Creep-, ——.o i | 1 -i - -, ___100 hours" 20 20 0 200 400 600 800100012001400160018002000 0 200 40 100012001400160018002000'" + Figure 53a Effect of Exposure Temperature and Figure 53b Effect of Exposure Temperature and Time On Difference In Room Temp- Amount of Prior Creep On Difference erature Tensile Properties Between In Room Temperature Tensile ProperAs-Exposed and Re-Machined Specimensties Between As-Exposed and Re-Machined of Rene 41 After Unstressed ExposureSpecimens of Rene' 41 After 10 Hours Exposure Exposure

0ZOO'F 1300'F 210 210 r z, tOO O o I a ~ )0. A Z OO 200 2 8 - - 4e + 0 A o190 190 0 - -- * I' _ I 20 180 180 "Deep racks" Code a 200 - - ------------ 200 - ~ 0*0 t As Treated 10 I0 h. o 0 I 4 I__~_. 5 h"r 190 ------ ------, ---- ----- ----- 190 ---- ----- ---- --- ^ - + V -- -* 4 + 0 hr Exposure 5A A 4 oohr Timte c 180 180 10OO hr 0. 025-inches Amount Remachined 170 he.0. 0. 091-inches from Exposed Diameter g 170 ---- ---- ---- ----- ---- ----- 170.2 />,,I I ) ~ A *, ~ l A 4 -40 Not Remachined -4, 160 1 -- 160 --. - - - + " 4' -- - "Deep Cracks", i.e. oxidized I 50 -- -- - --- ---- ----- ---- ---- 10 0/ ----- ---- ____ ____ ___ area on fracture surface 050 1 ~'~ L IS0 No thermocouples on gage O 150 t~ --- ----- ---- ---- ----- ---- 140 -/ ~ --- -------- -------- t —----- -------- -------- section on shoulders instead - No thermocouples on gage section - 93Q w^-,______ _______ _ _______ _________ __ ______ -_ —^_-_ ___g_^__ alumel wire tied around gage u1300 Xy,~,~~ -~- - section - oxidized area on fracture surface 120. 12].40 Thermocouples on gage section - tied on with chromel only 30 7 30 T A 4 Thermocouples on gage section - tied on with alumel only I "N ______"Normal" 20 ] i I Remachine 20 -- 0h0-0- a _itj' 10 \a'~'7^ --- -"Deep" o 0 W 0 - 10 --- ~. -— Not Remachined N i I I I,=-;Ao ~' IA: 0 "0 No thermocouples ~'I~\ -—....~I...on gage section 0.0. ~ — I 00 __ _____'.hi < - ~- ^ C\ I. Edge Cracking- - 0.10 1 0.10 - I> + ""7~~~~~~~~~~~~~~~~~ i ~~~~~~~~~"Deep" I:=aD.Oj~ I I I Desp~racks- 0.051 [ I I~.. Remachining. ---- Deep Cracks - 0.05 - -- --—./,_'' Normal Depth ofAchning Normalepth of Rmachinine o r h Figure 54 Effect of Prior Creep, Re-Machining, and Thermocouple Practice On Room Temperature Tenhsile Properties And Cracking Tendencies of Rene' 41 Exposed at 1200' and 1300' F 92

1400 Creep-Exposure 1600'F Creep-Exposure 1800'F Creep-Exposure 21l " 19019 0 F I - _ f 00 - -- / r 200 0 —-- -180 —--- 180 —--- c 190 —-------— 10- 170- 1 700-:_ I -18. 150 1 1801 3 —--- -150t --- 120 170 -.I~'O 160 - As Treated 0D 0 Or 4.5 hrs. Creep 50 / -- ---- -0 — --- - 120 -- -- 0 — 10.0 hrs. \ Exposure- ~ —-- --- ---- -----— 0 - -- 4 20.5 hrs. Time 0~1 J 1401! 10 Remachined- 110 --------------- wU.) I I /I \0 0. 025-in. from diam. 13 ----- --- --- -- - 100 0.002-in. from diam. 100 [i~ I Not Remachined c 120l''90 - -90 - 2i — -- - -- - -- 0 — - -- - -- - -- - 2 0- - -- -- -- -- o ——:-____ — oo*._ o___ _ — _ —_ _ o I T0~1 l I I I I I I 1 1 1 1 1, 11 3 3 21 Ir- Io - 20o,_ ~Uo 20;-_-_-~51 ~,-,- _` _ 1 1 _2 o0 1 2 3 4 5 6 7 8 0 1 3 4 o 4 5 6 7 8 9 10 0 I 2 3 4 5 6 7 8 0 1 Z 3 4 5 6 7 8 9 0 1 2 3 4 5 6 7 8 9 Total Plastic Strain - %Total Plastic Strain - % Total Plastic Strain - % Figure 55 Effect of Prior Creep And Re-Machining On Room Temperature Tensile Properties of Rene' 41 Exposed at 1400', 1600~, and 1800-F

Spec. R-146 Spec. R-127 Exposed 10 hours at Exposed 10 hours at 1200~F to 4. 70 percent 1300~F to 4. 30 percent deformation deformation (Not R e- Machined (Not Re-Machined Before Tensile Test) Before Tensile Test) a,;! \, 00 ) @$' 1 t $;: Spec. R-63 Spec. R-147 Exposed 10 hours at Exposed 97. 3 hours at 1300~F to 4. 60 percent 1300~F to 5. Z1 percent deformation deformation (Re - Machined (Not R e- Machined Before Tensile Test) Before Tensile Test) Spec. R-140 Exposed 10 hours at 1300~F to 4. 85 percent deformation Approx. 4x magnification or 1300F (Low Ductility and Oxidation on Fracture Surface for Subsequent Tensile Test at Room Temperature)

.07 07 0 12 F. Cracks indicated 1 300~F Stresses: 130,000 - 136,000 psi by dye test but Stresses: 108,000 - 114,000psi 045-inch ~_ ___ __ __ ___ _j(i.______ 0 i 0machined off before - Alumel wrapped o.06.606 tensile test nI / gage section I I~ / No~ thermocouple / on gage section.05 All couples tied / /.01'^^~ ^~ tpelt).|on with Chromel U O - D e' "e D Jol.. I ii o.04 - ---.04 All couples tied Fiur 57e Examples^^ on with Alumel. / o.U C.03.03 ubD'i: - -Initial Portion of Curves Omitted/ / -o for Clarity.02 1 1 j ^ -? Code (Applies to both temperatures)-.02 / - -- --— Spec. exhibiting deep cracking (i. e., low ductility and oxidation \,' All specimens with 3 thermoon fracture surface in room' it couples tied to gage section.01 --- -- ----- temp. tensile test). 01 ___ ^ I < —with Chromel or Alumel at....~ —- i.^-~' [ random - except as indicated -— Deep cracking not observed - "" 0 0 --- 0 I 2 3 4 5 6 7 8 9 10 11 0 1 2 3 4 5 6 7 8 9 10 11 12 13 Time - hours Time - hours Figure 57 Examples of Creep Curves of Rene' 41 Specimens Exhibiting Deep Cracking After 10 Hours Creep-Exposure at 1200~ or 1300~F.

Figure 58 Spec. No. R-90 (Section Showing ^,..'i 0 \ Surface) Exposed Without Stress ^' x >' " -l r* " -A.... 100 Hours at 1000~F,... -,. -,_ ~,.... _I -t-,^ ^ <. " - C*-RO o Subsequent Tensile Properties e)-~>^^N ^.... + i. Elongation: 22. 9% gK^1.^^ vrt ~Reduction of Area: 23. 2% /* cft' ^^ tCorresponding Re-Machined Spec. Y~^ i^<- ^Elongation: 24. 4%.... j -,t ->., Reduction of Area: 32. 4% ^^{^ ^ s.> T>-I ~ ~500x \ -.. s. -k 5O x Figure 59 Spec. No. R-99 (Section Showing Surface Attack) Creep-Exposure at 13000F for 100 Hours to 0. 63 C e[.~~'c Percent Deformation Subsequent Tensile Properties!~:~'~; Pk,',i'" < _Elongation: 17. 6% ~ {-~':'! _'',// Reduction of Area: 17. 2% 1 P< - < J } Approx. Properties of Re-Machined ~~~- I\~ \ )~ T~ SSpec. Elongation: 20% Reduction of Area: 24% -" 1 500x S x.i / /',. /....,....... S Figure 60 Spec. No. R-95 (Section Showing -^ _ n i - < Surface Attack and Isolated Crack)' t....;'/^3V'Creep-Exposure at 1300~F for 10 Hours to 2. 71 Percent Deformation "V.. i.^' ^f^..,,'C s Subsequent Tensile Properties 3y,'+> -. ***- (X;" ('"7 i^ +D"' Elongation: 10. 2% ~ /'*.... ) i V;- 1... Reduction of Area: 12. 9% yI i v s - * f I. t'/ S'",/ < " -s..' AN i' _ 1O O X )!I<~..,'','. —ti - g t / s: t.;r 5, 4 t, 96

Figure 61 Spec. No. R-93 (Section Showing I[r ~, Surface Attack) Exposed Without Co, o t o Q' ~ \ 1* ~ rt jStress 100 Hours at 1400~F > f5 ~/ s.. *. / Subsequent Tensile Properties'5, W 4 -,Elongation: 14.9% t # > b Reduction of Area: 14.9% - ] 1> <' ~ 4' e -:.? "<^.' Corresponding Re-Machined Spec. ~' ^^ ^ ^I'- Elongation: 22. 0%'' -.... 2. D Reduction of Area: 24. 9% f7 L 500x Figure 62 Spec. No. R-92 (Section Showing - -: r; r Surface Attack) Exposed Without;i4" Stress 100 Hours at 1600~F -^^^.,iK.'<.w1 ~ie;^;Subsequent Tensile Properties [c,.:~,..S. *Elongation: 6. 8% r'^ ^^^^^, ^ ^.*.....,~',)... Reduction of Area: 8. 3%:- ^p:.,, \'^Corresponding Re-Machined Spec. ".' {^~Ft~ <t~ t. -Elongation: 15. 1% __p.. -< -..~' Reduction of Area: 16. 6% Figure 63 Spec. No. R91 (Section Showing Surface Attack) Exposed Without Stress 100 Hours at 1800~F J'':~''1,~'EC~ 9 a ~8Subsequent Tensile Properties 7'- 5, ~ Elongation: 2. 8% Reduction of Area: 5. 0%,.>-;:,' n...-v-..::,, ~.,.:...s'c.: Corresponding Re-Machined Spec.,.',. _:..f.~';.i.. ~",' 00. Elongation: 4.2%' >Y.^ Reduction of Area: 5. 9%7'-"......:^ - -^.. 500x V.'.,. ~,..-,.~,, s;' i./,,,,,,~; i;.''.- ~ ~....: 3..*... 97

Figure 64 500x Figure 65 3500x Longitudinal Section of Spec. D-2 Electron Micrograph of Spec. D-Z Specimen D-2: Exposed Without Stress 474 Hours at 1700~F Figure 66 500x Figure 67 3500x Longitudinal Section of Spec. D-5 Electron Micrograph of Spec. D-5 Specimen D-5: Exposed Without Stress 2012 Hours at 1700~F 98 W ~ ~ ~ ~ ~ ~ ~ ~ ~ / A 0- ~ r~~~~~~~~~~~~~~~~~~~~~~~~~~ l 0 ~ JR A;:~4~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~..... *~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~~...

.01 ____________________ I 1800~F i.. ^^ 1 1 6006F U _ _ _ _ _ _ _____ 1I^ ^ 6 0 * 001 -— Code ^' <r ~Range of Iy u I Penetration, —4 U ^^ rJ) -,U) ~o ^/I.^13~00C~FI o l0F Penetration Greater Than ~~~~~~~~~~~~~~~~~~10001.0001 1 —----— /.00006 in. Detrimental /^' /~_~~~~~ ^^To Ductility L Limit of Detectability 00001 ------- 0. 1 10 100 1000 10,000 Exposure Time - hours Figure 68 Effect of Exposure Temperature and Time on Depth of Visible Surface Attack in Rene' 41

I 0I 1.0 rhQ' I- - - -i-;(D I _ _ I 01 *.1 5 - OI " —j.8 _ __ _ 4:. 4 -- -- ^ —- -10 hr 100 hr: Exposure Time o Q. 3 0 0 Elongation * 0.~2 A A Reduction of Area O -' -' *I- _,, I I t_, I I, I,,I i I 0.0005.0010.0015.0020.0025 Depth of Visible Surface Attack - inches Figure 69 Effect of Visible Surface Attack on Room Temperature Tensile Ductility of Rene' 41 Exposed Without Stress at 1300~- 18000 F

Figure 70 Section of Spec. D-Z Showing Attacked Surface Layers (Arrow shows solution of grain boundary phase in advance of general attack) Figure 71 2200x Electron Micrograph of "DepletedS" Layer of Spec. D-2 Spec. D-2 Exposed 474 Hours at 1700~F In Dilatometer Furnace 101

Figure 72 lOOOx Section of Spec. D-5 Showing Attacked Surface Layer (Exposed 1150 Hours at 1800~F In Dilatometer)

3. 608'' L e ( Matrix Lattice Layer 1: 3.5685 A 3. 604 Parameter in Layer 2:3. 5800 A 3. __ _ | SurfIndicated Layer Layer 3: 3. 5928 a 36. 001 1-.01 —-- e Mactrix: 35948 A 3. 596 \I,.j,. 3. 592 --—. - s - - ve.............-.-... 3.588 3. 584 ~ 3. 580 I-Oh' U h 3.576 ----- -1h rW X 3. 572 /! i 3.564 \i Surface Layer 1 Layer 2 Layer 3 Specimen Matrix 3. 560 oxide. 001 — -.2 00104-.0011 —-- 001 -4- - — Depth of Machining Principal 3. 556 subscale --- - depletion --— Overagedy' +carbides: Structural —.oxide- oxide needles I Feature 3 5 I I, I I I I Feature 0.001.002.003.004.005.006 Depth Below As-Exposed Surface - inches Figure 73 Effect of Surface Attack on Matrix Lattice Parameter of Spec. D-2 3.608 0-As Heat R: 1950~-1/2 hr+AC 3.604 Treated 1400~-16 hr+AC 3.600 - [R + 1700~ -474 hr] 2200~-15 hr+WQ 3.596 D 1500~- 100hr +AC _04 _D-2 Layer 3 D-2 Spec. Matrix-ce P 3.592 2 2- T —. —'.. - 2200'-15 hr+WQ ~I I __ 1500~-1000 hr+AC 3.588.. 1 -.. ——'. ---- 3+4Comparison of Compositions 584 -- jI' / \I // Rene' 41 D-2 Laer2 R v 3.580 / U Ht. R-134' Alloys of Ref. 3.576 1/ Curves of \' / Element wt. pct. wt. pct 3. 576-'[/ [ [' Beattie _ Cr 19:00 16-21 3.572 1 VerSnyder Co 11.00'10 3.56 -' - D-ZLayer I(Ref 20 ) C.09.12 d 3.568 — ~~"- Ti 3.2 2.5 s'7^ Al 1.6.7-1.0 3.564 1 Fe.3 Mn.07,1.25 5.556' Mo 9.75 0-11 Ni bal. bal. 5.552 - - - - - -- - - -, -— I —-—...... 0 1 2 3 4 5 6 7 8 9 10 11 Mo - wt. pct. Comparison With Data of Reference 20 Figure 74 Lattice Parameter Data for Specimen D-2 Compared to Data of Beattie and VerSnyder For Effect of Molybdenum Content on Lattice Parameter of Nickel-Base Alloys 103

I l c sIFUNCLASSIFIED UNCLASSIFIED THE UNIVERSITY OF MICHIGAN, Ann Arbor, Mich. THE UNIVERSITY OF MICHIGAN, Ann Arbor, Mich. EFFECT OF CREEP-EXPOSURE ON MECHANICAL PROP- EFFECT OF CREEP-EXPOSURE ON MECHANICAL PROPERTIES OF RENE' 41: Part II, Structural ERTIES OF RENE' 41: Part II, Structural Studies, Surface Effects, and Re-Heat Treat- Studies, Surface Effects, and Re-Heat Treatment, by Jeremy V. Gluck and James W. ment, by Jeremy V. Gluck and James W. Freeman, November 1961. 103p. incl. figs. Freeman, November 1961. 103p. incl. figs. tables and refs. (Project 7381; Task 73810) tables and refs. (Project 7381; Task 73810) (ASD TR 61-73, Pt. II) (Contract AF 33(616)- (ASD TR 61-73, Pt. II) (Contract AF 33(616)6462) 6I62) Unclassified report 6462) Unclassified report The effect of creep-exposure on room temp- The effect of creep-exposure on room temperature mechanical properties of Rene' 41 erature mechanical properties of Rene' 41 was studied for temperatures of 1200-1800~F was studied for temperatures of 1200-1800~F and times up to 200 hours. Unstressed ex- and times up to 200 hours. Unstressed exposures were for as long as 2012 hours at posures were for as long as 2012 hours at 17000F. Thermally-induced structural 1700oF. Thermally-induced structural changes reduced strength and ductility after changes reduced strength and ductility after exposures at 1400-1800~F. Reduced yield exposures at 1400-18000F. Reduced yield l|~ ~ l ~UNCLASSIFIED UNCLASSIFIED ( over ( over ) UNCLASSIFIED UNCLASSIFIED strength was due to decreased volume fraction strength was due to decreased volume fraction of gamma prime and secondarily to an in- of gamma prime and secondarily to an increase in the particle size. Ductility was crease in the particle size. Ductility was reduced by formation of massive grain bound- reduced by formation of massive grain boundary carbides. Up to 1500~F, creep caused ary carbides. Up to 1500~F, creep caused strain hardening and Bauschinger effects. strain hardening and Bauschinger effects. Except for surface effects, damage was re- Except for surface effects, damage was re- storable by re-heat treatment. Yield storable by re-heat treatment. Yield strength was restored by re-solution and re- strength was restored by re-solution and reaging to produce fine gamma prime. Complete aging to produce fine gamma prime. Complete Ire-solution of carbides was required to re- re-solution of carbides was required to restore ultimate strength and ductility. store ultimate strength and ductility. Microcracking was not observed. Creep in- Microcracking was not observed. Creep induced intergranular surface cracking at duced intergranular surface cracking at 1200-1300~F which reduced ductility. Sur- 1200-1300~F which reduced ductility. Surface effects for exposures above 1400~F were face effects for exposures above 1400~F were thermally induced. General principles were thermally induced. General principles were formulated for damage to properties of formulated for damage to properties of nickel-base alloys. nickel-base alloys. 7 -- - _ _ __UNCLASSIF IED _ UNCLASSIFIED I _ _ _ __ __ _ ~._ __. __ _ J _. _ _ _ __ __

|~I UmlUNCLASSIFIED UNCLASSIFIED THE UNIVERSITY OF MICHIGAN, Ann Arbor, Mich. THE UNIVERSITY OF MICHIGAN, Ann Arbor, Mich. EFFECT OF CREEP-EXPOSURE ON MECHANICAL PROP- EFFECT OF CREEP-EXPOSURE ON MECHANICAL PROPERTIES OF RENE' 41: Part II, Structural ERTIES OF RENE' 41: Part II, Structural Studies, Surface Effects, and Re-Heat Treat- Studies, Surface Effects, and Re-Heat Treatment, by Jeremy V. Gluck and James W. ment, by Jeremy V. Gluck and James W. Freeman, November 1961. 103p. incl. figs. Freeman, November 1961. 103p. incl. figs. tables and refs. (Project 7381; Task 73810) tables and refs. (Project 7381; Task 73810) (ASD TR 61-73, Pt. II) (Contract AF 33(616)- (ASD TR 61-73, Pt. II) (Contract AF 33(616)6462) 6462) Unclassified report Unclassified report The effect of creep-exposure on room temp- The effect of creep-exposure on room temperature mechanical properties of Rene' 41 erature mechanical properties of Rene' 41 was studied for temperatures of 1200-18000F was studied for temperatures of 1200-1800~F and times up to 200 hours. Unstressed ex- and times up to 200 hours. Unstressed exposures were for as long as 2012 hours at posures were for as long as 2012 hours at 17000F. Thermally-induced structural 17000F. Thermally-induced structural changes reduced strength and ductility after changes reduced strength and ductility after exposures at 1400-1800~F. Reduced yield exposures at 1400-1800~F. Reduced yield l|[~ l UNCLASSIFIED UNCLASSIFIED ( over )( over ) - -_ - -- -- -- _1- -_" _ _ _____ — -- -- 4 UNCLASSIFIED UNCLASSIFIED l strength was due to decreased volume fraction strength was due to decreased volume fraction of gamma prime and secondarily to an in- of gamma prime and secondarily to an increase in the particle size. Ductility was crease in the particle size. Ductility was reduced by formation of massive grain bound- reduced by formation of massive grain boundary carbides. Up to 1500~F, creep caused ary carbides. Up to 1500~F, creep caused strain hardening and Bauschinger effects. strain hardening and Bauschinger effects. Except for surface effects, damage was re- Except for surface effects, damage was restorable by re-heat treatment. Yield storable by re-heat treatment. Yield strength was restored by re-solution and re- strength was restored by re-solution and reaging to produce fine gamma prime. Complete aging to produce fine gamma prime. Complete Ire-solution of carbides was required to re- re-solution of carbides was required to restore ultimate strength and ductility. store ultimate strength and ductility. Microcracking was not observed. Creep in- Microcracking was not observed. Creep induced intergranular surfcce cracking at duced intergranular surface cracking at 1200-13000F which reduced ductility. Sur- 1200-1300~F which reduced ductility. Surface effects for exposures above 1400~F were face effects for exposures above 1400~F were thermally induced. General principles were thermally induced. General principles were formulated for damage to properties of formulated for damage to properties of nickel-base alloys. U nickel-base alloys. - _ _UN-CLASSIFIED _ _UNC LASS IF IED I 1 I _ _ _ _ I — - - $