RELAXATION OF RESIDUAL STRESSES IN A NICKEL-BASE SUPERALLOY

Transcription

1 CSA- y;. 1C,'f & RELAXATION OF RESIDUAL STRESSES IN A NICKEL-BASE SUPERALLOY.' DUE TO DISLOCATION CREEP D. Viereck*, D. Lohe**, O. Vohringer*** and E. Macherauch*** * ABB Kraftwerke AG, Mannheim, FRG ** Fachgebiet Werkstoffwissenschaften, Universitat - Gesamthochschule Paderborn, FRG ** Institut for Werkstoffkunde I, Universitat Karlsruhe (TH), FRG Abstract The thermal relaxation behaviour of macro and micro residual stresses in shot-peened sheet specimens of the nickel base alloy NiCr 22 Co 12 Mo 9 (Inconel 617) during different annealing times at 673 K and 1123 K was investigated with X-ray stress analysis and transmission electron microscopy. The assessment of the experimental findings shows that the residual stress relaxation is influenced by thermally activated slip and dynamic strain ageing at 673 K and by dynamic strain ageing, dislocation creep and precipitation processes at 873 K. At 1123 K the relaxation behaviour directly in the surface is caused by recrystallization processes and below the surface by dislocation creep processes which are controlled by viscous glide and precipitation hardening. Introduction Solid solution hardened nickel base superalloys are used as structural materials for combustion chambers of land base or aircraft gas turbines. Due to fabrication processes and surface cleaning treatments by blasting, local deformations occur near the surface of these complex structures. Thus ma,cro and micro residual stresses are produced across the combustion chamber wall. As these residual stresses may influence deformation behaviour and lifetime of high temperature components, a stress relief heat treatment must follow the fabrication process of combustion chambers. In contrast to many investigations on the residual stress relaxation behaviour of steels and various non-ferrous alloys (see e. g. (1,2)), there exists no knowledge about the residual stress relaxation _behaviour and the controlling mechanisms during stress relief heat treatment of high temperature alloys. The present paper deals with the thermal relaxation behaviour of macro and micro residual stresses of the nickel-base alloy NiCr 22 Co 12 Mo 9 (similar to InconeI617), which was studied in the course of an extensive research programme investigating the behaviour of combustion chamber materials under static, quasistatic and cyclic foading (3,6). Experimental Procedure The investigations were carried out on the nickel-base superalloy NiCr 22 Co 12 Mo 9 which was supplied in form of a sheet of 3.3 mm thickness with an average grain size of 95,.,. m by the Vereinigte Deutsche Metallwerke (VDM) under the trade mark Nicrofer 5520 Co. The chemical composition (wt.-%) of this alloy (11.7 Co, 21.8 Cr, 9 Mo, 1.13 Fe, 1.19 AI, 0,49 Ti, 0.15 Si, 0.07 C, 0.06 Mn, bal. Ni) is similar to Inconel 617.' The sheet material was solution heat treated and quenched by the manufacturer. Specimens with a size of 25 x 17 x 3.3 mm 3 were produced from the sheet and then shot-peened in an air blast shot-peening equipment (3). A cast steel shot S170 with a hardness of HRC was used at a pressure of 1.6 bar until a coverage of 98 % was achieved. For annealing times t a ' 90 min, the stress relief heat treatment was carried out in a fluidized bed furnace with nitrogen inert gas, and for t a > 90 min in a vacuum furnace (p =.10-5 bar). Residual stresses were determined with X-rays on {220}-interference '

2 planes according to the sin 2 1jJ-method using CrKex-radiation (3). Depth distributions of residual stresses and half width breadths of the X-ray interference lines were analysed by successive electrolytical removal of surface layers. Standard metallographic procedures were applied to prepare specimens for transmission electron microscopy (TEM-). More details concerning the experimental procedures are given elsewhere (3). Results Fig. 1a summarizes the depth distribution of the macro residual stresses generated by shot peening (as received state) and its alterations by annealing for 30, 90 and min at 673 K, 873 K and 1123 K. Fig. 1b represents the appertaining half width breadths of the X-ray interference lines as a measure of mean residual strains and micro residual stresses, respectively. The amount of the macro residual stresses directly increases from the surface (-600 N/mm 2 ) -to a maximum value in a depth of mm (-700 N/mm 2 ) and then continuously decreases. The half width breadth increases with decreasing distance from surface according to the shot-peening induced dislocation hardening. The upper diagram of Fig. 1a shows for Ta = 673 K that in a distance from surface' 0.1 mm the relaxation of macro 'residual stresses is much stronger during the first 3 minutes than afterwards. At 873 K (middle diagram Fig. 1a) it is remarkable that after 2000 min and min the amount of the macro residual stresses is higher near the surface than in a depth of mm. The macro residual stress relaxation in layers beneath the surface is similar to that occurring at 673 K. At 1123 K (lower diagram Fig.1.a) all annealing conditions show maximum amounts of macro residual stresses directly at the surface. The depth distributions T a =673 K t a : - = 3 min o = 90 min III,,-/. ;..." II, l:a = 2000 min,,/.0," II = min 'h III,... --".O"" _.:e"/:...:cs-..v, III"O!,C""'" 'o-e-. T a =1123 K as received state -800 L..--_--_--...-_--'""-_-.-.._----'" a, , ,-----,---..., -800 L..--_-J..--_--...-_--'""-_-0100 a , ,.--..., "" ,05 0,1 5 0,2 0,25 2,5 \ Ta =673 K ta - = 3 min 0 = 90 min 2,1 \ l:a = 2000 min " = min 1,7 "o. as received state 0it't:A_'*! 1,3 ib ----i 0,9 Vi 2,5 T a =873 K cv cv en 2.5\ 2,1 c, 1,7 -b cv 2,1 -'0.c.... -, "8 cv 1,7 as received state.c. I 1,3 I... "0 'j ---a \ d 0,9.c. C\ c'c _ -0"","" , T a =1123 K as received state 1,3 / _.. -o ---e-.. 0"' c II 0,05 0,1 05 0,2 Q25 Fig. 1a.: Macro residual stress distributions at 673 K, 873 K and 1123 K annealing for 3, 90, 2000 and min Fig. 1.b: Half width bredth distributions at 673 K, 873 K and 1123 K annealing for 3,90, 2000 and min

3 -.. _ '..,. --- " --_.- --'.. -' -= _._ : ""._-..;: -_: after 2000 min and min reveal nearly constant residual stress levels of about -50 N/mm 2. From the upper diagram of Fig. 1. b it can be seen that at 673 K the half width breadth is hardly influenced by the annealing treatment. The half width breadth di'stributions at 873 K (middle diagram Fig. 1.b) are characterized by a relatively weak annealing sensitivity with distinct plateaus in deeper surface layers. Annealing treatments at 1123 K (lower diagram Fig. 1.b) completely change the depth distributions of the half width breadth. The values measured near the surface are lower than in larger distance from surface. With increasing annealing time the position of these maxima and minima shifts to the interior of the specimen. Fig. 2 a, band c show representative dislocation structures attached to the half width breadth depth distributions for the as received shot-peened state and after annealing for 2000 min at 873 K and 90 min at 1123 K, respectively. The dislocation structures of the as received shot-peened state (Fig. 2a) and of the annealed state at 873 K (Fig. 2b) show no significant differences. Directly at the surface, a homogeneous dislocation structure is visible characterized by a hi.gh dislocation density and thin deformation twins with high frequency. In layers beneath the surface the microstructure changes to a more planar dislocation arrangement with distinct slip band structures and lower dislocation density. At 1123 K (Fig. 4) the change of the half width breadth distribution is also caused by a distinct change in the dislocation structure. In the region of minimum half width breadth, large recrystallized areas with very low dislocation density are visible, whereas in the region of the half width breadth maximum, only a small recrystallized area is detectable and in deeper surface layer, there is no recrystallized area at all.. Besides these alterations of dislocation structure, also the carbide structure and shape change between Fig. 2a, 2b and Fig. 2c. In the as,received shot-peened state (Fig. 2a) only some blocky primary intergranular carbides are visible. The same is walid after annealing at 873 K (Fig.2b). However, annealing at 1123 K (Fig.2c), effects near the surface a strong precipitation of secondary carbides in the former deformation twins and on the grain boundary, whereas in larger distances from surface a relatively homogeneous carbide distribution occurs. Discussion Residual stresses can be reduced or completely relaxed by supplying thermal energy. Thereby the elastic residual strain ES associated with the residual stress a rs via Hooke's law is converted into a microplastic strain E R by suitable deformation processes (2, 4). The transformation described by, the relationship ES = ars/e = -AE p (1 ) can, for example, be achieved by dislocation slip or creep in the residual stress field. During annealing up to min at 673 K the relaxation of macro residual stresses is combined with microplastic strains E p, 0,2 %. Annealing for 3 min produces a relatively large relaxation effect compared to annealing times of 90 min or more (see upper diagram Fig. 1a). However, the micro residual stress relaxation, which is characterized in this investigation by the half width breadths of the X-ray interference lines, is insignificant (Fig. 1b). An assessment of this behaviour is possible by taking into consideration results of tensile tests and stress relaxation tests which were carried out on the same material at 673 K (3,5). In these cases, the deformation behaviour is influenced by the elastic interaction of diffusing interstitial foreign atoms (carbon and nitrogen) with slip dislocations. This well-known dynamic strain ageing process is accompanied by a strong pinning of slip dislocations by the solute atoms, resulting in an impediment of stress relaxation (3). Therefore it can be assumed that the relaxation of macro residual stresses at the beginning of annealing at 673 K is caused mainly by the thermally activated slip of dislocations interacting with short range obstacles (3, 5). However, with increasing annealing time, the pinning effect intensifies and strongly reduces the relaxation rate. Due to the directed dislocation slip in the residual stress field, only a small rearrangement of the dislocation structure and a negligible annihilation rate of dislocations occur (3). For this reason, the insignificant relaxation of micro residual stresses at 673 K is understandable. -; '/,

4 Fig. 2: Dislocation structures attached to the half width breadth depth distributions of 2,9 V; cu cu 2,5 cu :s +- "'8 cu "C d 2,1 1,7 1,3 0,9 0 0,05 0,1 0,15 0,2 distance from surface (mm] a) the as received 0,25 shot-peened state, 3,3 2,9r i' L,J... :s i i L l j l ] :,J 0, T =873 K t =2000 min o, o 0,05 0,1 0,15 0,2 0, _.. _... b) annealed state at Ta = 873 K, t a = 2000 min, V; cu C' 3 1,7 +- "'8..0.&:. +- "C 1, d / o/"o 0,9 0 0,05 0,1 0,15 T =1123 K t = 90 min 0,2 c) annealed state at 0,25 Ta = 1123 K, ta = 90 min

5 -."'.,..". -_'" ' _".- _ -..,.- _..,,' _ _ - -, _-. ' '" -' '" '.. Mechanical tests at 873 K show that both dynamic strain ageing and creep processes determine the deformation behaviour of this material (3,5,6). From creep tests, at this temperature the evaluation of the stress dependency of the secondary creep rates by the Norton-law Es = A an (2) yields to the creep exponent n = 6. To get more information about the controlling mechanisms in the case of residual stress relaxation, mean plastic strain rates were determined from the depth distributions of macro residual stresses in Fig. 1a (middle diagram). As schematically shown in Fig. 3, A ars-values at adjacent annealing times At a = t2 - t1 (t2 > t1) at depths s = canst. yield together with eq. (1) the strain rates E = A p / At = -A s / At = _Aa rs / (EAt a ). (3) en to. \!) -+- d :::J :g c.. 0 t.""..,. 2 -r--- tj.g rs distance from surface s ---- Fig. 3: Schematic plot of residual stress vs. distance from surface for evaluation of mean strain rates during relaxation. Plots of mean values of.e and ars at s = canst. and anne.aung temperatures Ta = 873 K and 1123 K are given in Fig.4 in log E - logars-diagrams. The slopes of the regression lines are identical with the Norton exponent n. At 873 K 10-4 T a =873 K LJ6 c s=o,05m// /' 5=0,1 mm (upper diagram Fig. 4), large n-values at depths s = 0.05 mm and 0,1 mm are evaluated at high residual stresses and large strain rates (which means short annealing times). In this area power-low breakdown occurs, and relaxation of macro residual stresses is impeded by dynamic strain ageing. At smaller residual stresses, a n-value of 6 is determined, which indicates that climbing of originally pinned edge dislocations initiates a diffusion controlled dislocation creep process, the so-called viscous glide. The relatively high amounts of macro residual stresses directly at the surface after annealing for 2000 min and min at 873K (Fig. 1a) can be attributed to thin Si0 2 areas formed during the annealing treatment and detected microanalytically. These areas are characterized by a smaller thermal expansion coefficient than that of the metal matrix. They produce, after annealing and coolin.g to room temperature, increased compressive residual stresses in both the 5i02 and the metal matrix areas at the very surface residual stress [N/mm 2 ] Fig. 4: Stress dependency of mean strain rates determined from macro residual stress relaxation at Ta = 873 K and Ta = 1123 K at distances from surface s = canst.

6 2,5 T a =873 K 200 t a =3 min 150 1,7 '0 HWB 100,/ 1,3 'c 0,9 N I Vi 2,5 T a =873 K 200 E u c- t a =90 min O' C\ :s 2,1 150 os:..r::. "'8 0v; 1,7 /HWS 100 C c-.c ""0 o-a-a r::. "'0 'j 'r- \ c 1,3 SO 0 :.;: '6-0 c;..r::. 0,9 0 -vi :.a 2,5 T a =873 K 200 t a =2000 min 2,1 150 HWB 1, t 1,3 " SO -----t 0,9 0 0,05 0,1 0,15 0,2 0,25 0 (j U At 873 K the relaxation of macro residual stresses is accompanied by decreasing half width breadths and micro residual stresses, respectively (Fig. 1b). However, the micro residual stresses relax slower than macro residual stresses. Examples of the relaxation behaviour of half width breadths are presented (in an other scale) in Fig. 5 together with dislocation densities determined by TEM. Obviously, due to the larger driving force, annealing treatments at 873 K cause a larger annihilation rate of dislocations at small distances from surface than at higher s-values. This behaviour is responsible for the formation of distinct plateaus in the HWB- and Pt-depth distributions, which can also be seen in Fig. 5 in the depth region of the plateaus. The dislocation rearra,ngement and annihilation during the relaxation process may be impeded by the precipitation of extremely fine carbides on slip dislocations. Near the surface precipitation of carbides occurs, too, but mainly at the boundaries of the numerous deformation twins. Therefore, it may be possible that sufficient mobile dislocations are available for further relaxation of micro residual stresses at small s-values. Fig. 5: Half width breadth and dislocation density distributions after ann,ealilng at Ta = 873 K for t a = 3, 90 and 2000 min Annealing of shot peened states at 1123 K yields residual stress relaxation which is at low s-values mainly a result of recrystallization processes and at larger s-values a result of creep controlled dislocation rearrangements (Fig. 2c). The evaluation of the Norton exponent n = 6 by the aid of eq. (3) shows that dislocation creep by viscous glide determines the macro residual stress relaxation in depths between 0.05 and 0.1 mm (lower diagram Fig. 4). At s = 0,25 mm the same mechanism is effective for relatively short annealing times (large e and errs values). However, at longer annealing times (small E and ars-values) a transition to n = 11 occurs which is caused by carbide precipitation on dislocations. Consequently, retardation of residual stress relaxation observed in this distance from surface is caused by the pinning of mobile dislocations by carbides (lower diagram Fig. 1a). Acknowledgement The financial support for these investigations by the Deutsche Forschungsgemein.schaft is gratefully acknowledged. They were performed within the IISonderforschungsbereich 167". References 1. J. Hoffmann, Dr.-Ing. thesis, Universitat Karlsruhe (1985) 2. O. Vohringer, in: Internat. Guidebook on Residual Stresses; Advances in Surface Treatments, Vol. 4, p. 367, Pergamon Press (1987) 3. D. Viereck, Dr.-Ing. thesis, Universitat Karlsruhe (1989); Fortschritt-Berichte VDI, Reihe 5, Nr. 202, VDI Verlag Dusseldorf (1990) 4. M. R. James, in: Internat. Guidebook on Residual Stresses; Advances in Surface Treatments, Vol. 4, p. 349, Pergamon Press (1987) 5. D. Viereck, G. Merckling, K. H. Lang, D. Eifler and D. Lohe, in: Festigkeit und Verformung bei hoher Temperatur, p. 201, DGM-Informationsgesellschaft, Oberursel (1989) 6. G. Merckling, Dr.-Ing. thesis, Universitat Karlsruhe (1989),