Thermal Decay of Shot-Peening Induced Residual Stresses During Annealing of a Bake-Hardening Steel

Thermal decay of shot-peening induced residual stresses during annealing of a bake-hardening steel

A. Rossler, J.K. Gregory

Technische Universitat Munchen D-85747 Garching, Germany EMail: roessler(a)w. mw. turn, de

Abstract

The bake hardening steel ZSTE 220 BH was shot peened to Almen intensities of 0.12, 0.17 and 0.27 mmA. Residual stress profiles were measured using x-ray diffraction. Specimens shot peened with 0.17 mmA were subsequently annealed at temperatures of up to 250 °C and the change in residual stress profiles was measured. While thermal residual stress decay was high at the surface, residual stresses below the surface were stable. This result is explained by invoking the strain aging effect as investigated in the tensile test. For a low degree of cold work, a pronounced strain aging effect stabilizes residual stresses at 170°C, while a high degree of cold work tends to eliminate the strain aging effect.

1 Introduction

Mechanical surface treatments such as shot peening or deep rolling are routinely applied to components in machinery with the aim of improving fatigue performance. Particularly in high strength materials, this improvement can be attributed to the presence of compressive residual stresses in the near surface layer which retard and may even arrest crack growth. Hence, conventional wisdom has thus far been to avoid any thermal treatment after shot peening, because this is expected to cause thermal decay (i.e. reduction) in the magnitude of the residual stresses. Indeed, most of the available literature which discusses shot peening followed by annealing has concentrated on the aspect of thermal decay, e.g.,(V6hringer [1]).

312 Surface Treatment

Two compelling reasons exist for re-examining the question of thermal treatment following shot peening. Firstly, shot peening plus annealing has been shown to yield superior fatigue properties for certain age-hardenable alloys (Fair [2], Gregory [3], Wagner [4]). This effect has been attributed to a retardation of thermal decay of residual stresses owing to simultaneous age-hardening (Berger [5]). Indications are available that annealing after shot peening can also be beneficial in carbon steels, provided that the time and temperature are judiciously chosen. For a high strength steel tempered at 300 °C following shot peening, this effect was attributed to stabilization of the dislocation networks which give rise to the compressive residual stresses (Qiong [6]). These thermally stabilized networks are then more resistant to degradation by cyclic loading than after shot peening alone. Secondly, the application of coatings for corrosion protection (i.e. paint) after a mechanical surface treatment generally involves a temperature excursion of up to 150-180 °C. These temperatures correspond to those associated with strain aging effects by the formation of Cottrell atmospheres, i.e, carbon and/or nitrogen atoms in the matrix diffuse to dislocations, resulting in pinning effects which increase strength. The present work aims to determine the effect of annealing in this temperature range on residual stress profiles in a bake hardening steel which can be expected to exhibit particularly pronounced strain aging effects.

2 Material and Experimental Procedure

The bake-hardening steel ZSTE 220 BH was received as thin plate with the composition shown in Table 1. The material had been pre-strained to between

1 and 1.5%. The average grain size was 15 um and grains were slightly elongated in the rolling direction. The as-received pre-strained condition consists primarily of ferrite grains with isolated pearlitic regions. Grain- boundary cementite is also present.

Table 1. Chemical composition of ZSTE 220 BH (wt.%)

P Al Elements C Si Mn S N ZSTE 220 BH 0.038 0.04 0.25^ 0.024 0.013 0.0028 L 0.043 J

The bake-hardening (BH) -effect was determined in tensile tests on standard specimens oriented transverse to the rolling direction. The strain rate was kept constant at 4 x 10'Vs. The BH-effect was quantified as the difference between the lower yield point in material after strain aging and the 0.2% offset in material in the as-received condition. In cases where the material was plastically pre-strained, the BH-effect was taken as the difference between the lower yield point after strain aging and the stress reached immediately before unloading. Care was taken to calculate the stress using the relevant specimen cross-section, i.e., that obtained just before unloading. The objective of these tests was to determine the influence of the annealing parameters (T, t) and pre-strain on the

Surface Treatment 313

BH-effect. In order to determine the influence of high values of pre-sii am on the BH-effect, material was cold-rolled 50 %. Heat treatments were carried out between 140 and 250 °C in a silicone oil bath with temperature controlled to within +/- 1°C. Additional tensile tests were carried out at elevated temperatures comparable to the annealing temperatures. Flat specimens with dimensions 30 mm x 12 mm x 1.7 mm were shot peened on one side with SI 10 shot using the parameters shown in Table 2. The zinc coating present on the sheet was electrolytically removed prior to shot peening.

Table 2. Shot peening parameters

Intensity Pressure Coverage Mass flow [mmA] fbar| [%I [kg/mini SP 1 0.12 0.7 98 0.5

SP2 0.17 1.0 160 1.5 SP3 0.27 2.0 250 1.5

Residual stresses were measured by x-ray diffraction using the snr-ij/ method. Lattice strains were determined on the {211} planes in ct-Fe with CrKct- radiation. Stresses were calculated from the strains using the x-ray elastic constants 1/282 = 5.81 x 10~* (MPa)"', Si= -1.27 x 10"* (MPa)"\ Depth profiles were measured by electropolishing to successively remove surface layers.

Stresses were corrected according to (Moore [7]). Thermal stress relaxation was measured using one single specimen per depth by first electropolishing to the depth of interest and then determining the residual stress before and after annealing. Possible variations in absolute values between different specimens could thus be excluded.

3 Results

The influence of various annealing parameters as well as test temperatures is shown in Fig. 1. The as-received condition exhibits a yield stress of 251 MPa, after annealing 30 min at 170 °C a lower yield point of 279 MPa is found. This difference is designated the BH-effect. Testing at slightly elevated temperatures after annealing results higher yield stresses than that of the as-received condition. The Liiders strain increases strongly with increasing annealing time and temperature and achieves values of up to 6 %. Beyond the Luders strain, the stress-strain curve merges with that of the as-received condition, as can be seen in Fig. 1. Fig. 2 shows the significant influence of time and temperature on the BH-effect for a pre-strain of 3.5 %. Fig. 3 shows both the ultimate tensile strength and the BH-effect as a function of annealing time at 170 °C for a pre-strain of 2.5 %. A correlation between the ultimate tensile strength and BH-effect is evident. Tensile properties after cold rolling 50 % were ao.2=550 MPa, auis"650 MPa with an elongation of less than 5 %. No BH-effect could be found by annealing this condition.

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350 -i annealed 30 min 170°C, tested at RT 325 - I annealed 120 inin 200°C, tested at 200°C I as-received CL, 300 -

^2754

% 250 ,'f annealed 2000 min MOT, tested at 140°C I \ I " annealed 240 mm 170°C, tested at 170°C j 200 ^-J 0 1

strain [%]

Fig. 1 Effect of different annealing and test temperatures on stress-strain

behavior

30 -[• 140°C

100°C

0 20 40 60 80 100 120 140 160 annealing time [min]

Fig. 2 Effect of annealing temperature on BH-effect for a pre-strain of 3.5%

ultimate tensile strength -*- BH-effect 20 100 1000 10000 100000 annealing time [min]

Fig. 3 Ultimate tensile stress and BH-effect for a pre-strain of 2.5% at 170°C

Surface Treatment 315

Shot peening causes the surface layer to work harden, the degree of work hardening increasing with Almen intensity. This can be assessed both by the full peak width of the diffraction line at half-maximum intensity (FWHM, Fig. 4) and by the microhardness. For steels, the tensile strength can be estimated from the microhardness. This estimated is plotted together with the residual stress profile in Fig. 5 for SP2. Both GI and

2.0

SP 2 1=0.17 mmA SP 3 1=0.27 mmA SP 1 1=0.12 mmA T 1.5

0.54- 0 mo 200 sm 400 500 distance from the surface

Fig. 4 FWHM induced by different shot peening parameters

- 400 t:

^350 -300

distance from the surface (urn)

Fig. 5 Residual stresses and local tensile strength as calculated from

microhardness for SP 2

316 Surface Treatment

as shot peened annealed 20 min at 170°C

annealed 20 min at 250°C

-400 -L distance from the surface

Fig. 6 Residual stresses after shot peening treatment SP2 and annealing at different temperatures

4 Discussion

4.1 Bake hardening effect

Increasing annealing temperature leads to a more pronounced BH-effect owing to the more rapid diffusion of interstitial atoms. Depending on pre-strain, a saturation value of 40-50 MPa is reached for sufficiently long annealing times (Fig. 7).

mo iso annealing time [minj

Fig. 7 Effect of pre-strain on BH-effect by annealing at 170°C

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The development of the BH-effect with time can be divided into two regimes. For annealing times t < 80 min and T= 170 °C, pre-straining gives rise to a strong BH-effect. The increase in dislocation density associated with a strain of at least 1.5 % is sufficient to cause significant strain aging within a few minutes, as can be seen in Fig. 7. For t > 80 min, pre-straining tends to reduce the BH- effect as annealing time is increased. The higher the dislocation density, the smaller is the BH-effect. The observation that the BH-effect is absent in material cold-rolled 50 % is consistent both with these results and with the results in (Elsen [8], Engel [9]). No clear picture is available from the literature regarding the effect of pre-strain on the BH-effect (Elsen [8],Engel [9]). This is presumably related to the variation in thermomechanical treatment. In some cases, discrepancies can be attributed to an improper choice of specimen cross- section when calculating the stress. At a pre-strain of 25 %, (Engel [9]) found the achievable BH-effect to be reduced by 50 %. Based on these results, it can be said that the BH-effect can be exploited in reasonable annealing times by proper choice of annealing parameters for a certain minimum plastic deformation. As demonstrated in this work, a pre-strain of 1.5 % is sufficient to activate the BH-effect for annealing times of several minutes. However, if the degree of plastic deformation is too high, the BH-effect is reduced. For a high dislocation density, the influence of Cottrell atmospheres developed by strain aging on an increase in yield point is minor, the major strengthening effect being strain hardening itself. For annealing times t > 5000 min at T=170 °C, a reduction in the BH-effect can be detected. This can be attributed to an overaging effect. The concentration of interstitial carbon and nitrogen atoms at dislocations reaches a saturation value, above which carbides or nitrides form. The strengthening mechanism from these fine particles is not as effective as for the Cottrell atmospheres. This is supported by the similar dependence of the BH-effect and GUTS on annealing time as shown in Fig. 3.

A remarkable result is the fact that higher test temperatures do not cause the yield stress to decrease (Fig. 1). This is presumably related to a dynamic strain aging effect. At these temperatures, the resistance to plastic deformation decreases relative to that of material annealed and then tested at room temperature, but is still at least as high as that of shot peened material.

4.2 Thermal stress relaxation

As expected, the greatest thermal decay in compressive residual stress was found at the surface. Both long range and short range residual stresses relax with increasing annealing temperature. According to (Vohringer [1]), long range residual stresses can be relieved by dislocation motion, while the relief of short range residual stresses requires rearrangement and annihilation of dislocations.

Fig. 8 illustrates the fact that no significant reduction in short range residual stresses at a depth of 50 um is found after 20 min at 170 or 250 °C.

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1.6

1.4 - as shot peened

annealed 20 min at 170°C 1.2 - annealed 20 min at 250°C

0.8 4

0.6 ---I

50 WO 150 200 250 300 distance from the surface [jum]

Fig. 8 FWHM for different annealing temperatures

Fig. 6 shows the residual stress profiles for the same annealing parameters. The long range residual stresses remain constant beyond a depth which depends on annealing time, as do the short range residual stresses. The rapid thermal stress decay at the surface can be explained by the high dislocation density in these regions. When thermal energy is supplied, dislocation motion is immediate, leading to enormous reductions both in long range and short range residual stresses. This can be seen in Fig. 6 by comparing thermal decay at 0 and

35 um. According to (Vohringer [10]), thermal relaxation of residual stresses occurs when QRS > Oyc/t, where

170 °C. This characteristic depth shifts to higher values for an anneal at 250 °C. As demonstrated by the microhardness results, the degree of cold work decreases with depth. The BH-effect becomes correspondingly more significant as the dislocation density decreases. The formation of Cottrell atmospheres associated with strain aging pins the dislocations (Argon [11], Dahl [12]). This phenomenon increases resistance to plastic deformation and hence to thermal stress decay. Usually, the yield stress decreases significantly with increasing temperature, as was found for microalloyed steels when temperature is increased from 25 to 200 °C (Burgahn [13]). Similarly, cjyc/t also decreases with increasing annealing temperature and time (Vohringer [10]). For shot peened material, the value of cjyc/t is a function of depth, owing to the microstructural changes

Surface Treatment 319

induced by shot peening itself. All barriers to dislocation motion which are thermally stable are expected to retard thermal stress decay and also to shift the degree onset of decay to higher temperatures.

Assuming that cjyc/t exhibits a temperature dependence comparable to the yield stress, these observations are consistent with the concept that ayc/t is the parameter relevant for thermal stress relaxation. The thermal stability of the yield stress in BH-steels in the temperature range of 170 °C is thus related to the formation of Cottrell atmospheres. Two effects can be considered to be particularly relevant: - resistance to plastic deformation decreases with increasing temperature

owing to thermally activated dislocation motion - the potential BH-effect achievable by annealing depends on the degree of deformation. The interplay between these effects influences dislocation motion, which is responsible for thermal stress decay via ayc/t.

5 Summary

The BH-effect as determined in a tensile test depends strongly on annealing time and temperature. The maximum BH-effect is achieved for a pre-strain of 1.5 % and annealing 30-40 min at 170 °C.

The development of Cottrell atmospheres increases resistance to plastic deformation. Even at testing temperatures of 170 °C, this resistance is high compared to that of the as-received condition.

The BH-effect is absent for high degrees of cold work. BH-effects, i.e., strain aging can be invoked to rationalized the thermal stress relaxation behavior. The beneficial effect of strain aging in retarding residual stress relaxation can be explained assuming that a^, which determines the degree of relaxation, exhibits a temperature dependence similar to that of the yield stress. At the surface, where the dislocation is high, no BH-effect is possible and the usual processes which lead thermal decay take place. Well below the surface, the BH-effect increases the resistance to plastic flow and residual stresses are stable with respect to thermal decay.

References

[1] Vohringer, O., Relaxation of Residual Stresses by Annealing or Mechanical Treatment, Advances in Surface Treatments A. Niku-Lari Vol 4 Pergamon Press, pp. 367-396, 1987.

[2] Fair, G., & Noble, B., & Waterhouse R.B., The Stability of Compressive Stresses Induced by Shot Peening under Conditions of Fatigue and Fretting Fatigue, Advances in Surface Treatments, (A. Niku-Lari) Vol. I, Pergamon Press, pp. 3-8, 1984.

320 Surface Treatment

[3] Gregory, J.K., & Miiller, C, & Wagner, L., Preferential Surface Aging: Novel Methods for Improving the Fatigue Life of Mechanically Loaded

Components (in German), Metall 47, pp. 915-919, 1993.

[4] Wagner, L., & Gregory, J.K., Improve the Fatigue Life of Titanium Alloys - Part II, Advanced Materials and Processes, 145, pp. 50HH-50JJ, 1994.

[5] Berger, M.-C., & Gregory, J.K., Residual Stress Relaxation in Shot Peened TIMETAL 21s, Mafer. &%. Engg. /4 263/2, pp. 200-204, 1999.

[6] Qiong, Q., & Renzhi, W., Influence of the Change of Microstructure and Residual Stress Field in the Surface Shot-Peening Straining Layer on Fatigue Behavior, Shot Peening: Science * Technology * Application

(Wohlfahrt, H., Kopp, R., Vohringer, O.), DGM Oberursel, pp. 231-238, 1987.

[7] Moore, M.G., & Evans, W P., Mathematical Correction for Stress Layers in x-Ray Diffraction Residual Stress Analysis, Trans. SAE 66, 1958.

[8] Elsen, P., Bake-Hardening-Effect in Thin Sheet (in German) VDI Forschungsberichte Reihe 5 Nr. 314, 1993.

[9] Engel, B., & Drewes, E.J., Investigations of the Buckling Behavior of High Strength Thin Sheet with Respect to Strengthening due to Thermal Influence from the Varnishing Process (in German), Stahl und Eisen 103 #r. 77, pp. 819-824, 1983.

[10] Vohringer, O., & Wohlfahrt, H., Relaxation of Residual Stresses (in German), Eigenspannungen und Lastspannungen HTM, pp. 144-155, 1976.

[1 1] Argon, A.S., & Ashby, M.F., Thermodynamics and Kinetics of Slip, Progress in Material Science, Volume 19, 1975.

[12] Dahl, W., Principles of Strength, Toughness and Fracture (in German),

7, 1983.

[13] Burgahn, F., Unidirectional Deformation Behavior and Microstructure of selected Steels as a Function of Temperature and Strain Rate (in German),

Dr.- Ing. Diss. TH Karlsruhe, 1991.