M. Pahutová

Dislocation structure and applied, effective and internal stress in high-temperature creep of alpha iron

Abstract The dislocation structure of alpha iron formed in high-temperature creep mas investigated in the transmission electron microscope and its quantitative characteristics—dislocation density, subgrain diameter and subgrain misorientation—were measured in the primary as well as in the steady-state creep. The steady state substructure characteristics are correlated with the steady state values of the applied, effective and internal stress and the development of the substructure in the primary creep is discussed. The results of the structure observation correspond well to the macroscopic creep data (Pahutora, Orlova, Kuchařova, Cadek, to be published).

High temperature creep in copper

Creep in copper of 99·99% purity in the temperature interval 550-1025°K was investigated by the isothermal tests technique. Two dislocation mechanisms operating in parallel were detected. In the region of lower creep rates-Region 1-the apparent activation energy of creep depends linearly on stress: [image omitted] where Q01 = 47·4 ± 2·8 kcal mol-1 and B1 = 3·2 ± 0·5 kcal mol-1 Kg-1 mm2. The value of Q01 is close to the activation enthalpy of lattice self-diffusion in copper. Consequently, creep is controlled by lattice self-diffusion. The stress exponent n ≡ n1 = δ In ∊/δ In σ is a function of both stress and temperature. Generally, n1 decreases, reaches a minimum and increases again with the increasing stress. The lowest value of n1 ≃ 5·4. Non-conservative motion of jogs on screw dislocations dependent on lattice self-diffusion was suggested to be a rate-controlling process in Region 1.In the region of higher creep rates-Region 2-the apparent activation energy Qc2, is considerably higher than the activat...

High temperature creep of alpha iron

Abstract The creep of decarburized ARMCO iron has been, studied in the temperature range of 820–1170 °K by the isothermal tests technique. In the range of 820–920°K the apparent activation energy of creep depends linearly on stress; the value corresponding to the stress σ= 0, H0 = 105.0 ± 1.9 kcal mol−1, is by a factor of about 1.75 higher than the activation enthalpy of lattice self-diffusion H sd . In the range of 920°K-TC (TC = 1042 °K, the Curie temperature) the apparent activation energy depends both on stress and temperature, achieving the value of 220 kcal mol−1 at TC. Accepting the model of nonconservative motion of jogs on screw dislocations, the high value of H0 is interpreted on the basis of a suggestion that the nonconservative jump of a jog on a screw dislocation by one interatomic distance is connected with co-operative absorption or emission of an effective number of m vacancies, so that H 0 = m H sd . The temperature dependence of H0 in the range of 920 °K-TC is interpreted by the temperature dependence of H sd in this temperature range. If the power function of σn type is chosen to describe the experimentally determined stress dependence of the minimum creep rate, n depends both on stress and temperature; at a given temperature it decreases initially with the increasing stress, achieves a minimum and then increases again. The experimentally determined dependence n = n(σ, T) is correlated with the dependence following from the accepted model of nonconservative motion of jogs on screw dislocations under the assumption that the ratio of the density of moving screw dislocations ρs, to the density of dislocations unbound in subboundaries ρ is independent on stress.

Some stress change experiments on creep in α zirconium

Strain behaviour after various applied stress changes in the creep of zirconium at 873 K is described and discussed. After reduction of the applied stress σbyΔσ ≈ σ time-dependent backward strain ϵb occurs, amounting to almost three times the elastic strain ϵe that corresponds to the stress reduction. The ratio ϵb/ϵe decreases with applied stress. After some time the backward strain is replaced by forward strain ϵf; the magnitude of the forward strain amounts to about 1/5 of the maximum backward strain ϵbmax. No incubation period was observed after stress reductions Δσ < ΔσcwhereΔσc is the stress reduction to which zero creep rate (as measured immediately after the stress reduction) corresponds, and which is identified with the mean effective stress σ∗. No incubation period was observed after reductions in stress σ1byΔσ ≈ σ1 for a time period tr and application of a stress σ2 < σ1. Even at short time periods tr and low stresses σ2 the initial strain rate was positive. When the forward strain had reached a maximum value, however, backward strain occurred; it was replaced — after a time — by the “second” forward strain. The forward strain rate finally reaches the steady state value corresponding to σ2. Discussion of results is based on the network growth model originally suggested by Mitra and McLean and, with regard to the behaviour of sub-boundaries in creep, the results of Exell and Warrington and Myshlyaev et al.

High temperature creep of a Cu-16 at.% Al alloy

Abstract The creep of a Cu-16 at.% Al alloy in the temperature range of 670–1100°K has been investigated by the isothermal tests technique. Two dislocation mechanisms operating in paralell were observed in the given temperature range. In the region of lower temperatures ( 10 4 T ⪆ 11.5 ) the apparent activation energy Q r 1 depends linearly on stress. The value of Q 01 corresponding to the stress equal zero is close to the expected activation enthalpy of lattice self-diffusion of copper in the Cu-16 at.% Al solid solution. The stress dependence of minimum creep rate can be described by the power function of the σ n type, where n = 6.25. The results for region 1 can be correlated with the model of nonconservative motion of jogs on screw dislocations dependent on lattice self-diffusion. In the region of higher temperatures (region 2) the mechanism dominates characterized by the activation energy that is higher than the expected value of the activation enthalpy of lattice self-diffusion of any component of the investigated solid solution. The stress dependence of creep rate can be described by the power function of σ n type, where n = 5.40. The dislocation mechanism dominating in region 2 was not identified.

Steady-state creep in alpha iron as described in terms of effective stress and dislocation dynamics

Abstract Using the strain transient dip test technique the mean effective stress [sgrave]∗ was measured in steady-state creep of alpha iron in a temperature interval 400 to 700°C. At temperatures 400 to 450°C the mean effective stress was found to increase with increasing temperature, while at temperatures 550 to 700°C it is temperature-independent. At intermediate temperatures a more complicated behaviour of mean effective stress was observed. The parameter m∗ = (∂ ln v g/∂ ln [sgrave]∗) describing the effective stress sensitivity of dislocation glide velocity v g was estimated, assuming the moving dislocation density ρ m to be a constant fraction of the density ρ of dislocations unbound in sub-boundaries. The contribution to the effective stress dependence of steady-state creep rate of the effective stress dependenoe of ρm was found to be considerably smaller—at least at high effective stresses—than that of v g. In the temperature interval 550 to 700°C the apparent activation energy of dislocation glide...

Transients in the creep of a 16Cr12Ni2.5Mo austenitic steel II: Structure

Abstract Using transmission electron microscopy the changes in dislocation structure of a 16Cr12Ni2.5Mo austenitic steel were investigated during transient creep after an applied stress reduction from σ1 to σ2 = 100 MPa at temperaturesof 973 and 1023 K. The stress reductions were performed after a creep strain of 0.23 had been reached; this is well within the steady state region at σ1 = 300 MPa at 973 K and σ1 = 200 MPa at 1023 K. It was found that the formation of a new substructure during transient creep after a stress reduction from σ1 to σ2 is preceded by the disintegration of a significant fraction of the subboundaries formed during creep up to the steady state at the stress σ1. The subboundary disintegration is related to the glide of dislocations that had been bound in these subboundaries. Consequently, the density ϱm of mobile dislocationsincreases despite the fact that the density of ϱ of free dislocations does not change significantly. This makes it possible to account for the observed non-monotonic increase in the forward creep strain, i.e. the point of inflection on the transient creep curve after stress reduction.

Some basic creep characteristics of ZrSn&.z.sbnd;Mo and ZrSnMoNb alloys part I. Steady state creep

Abstract The steady state creep rate, dot ϵ s , of five ZrSnMo alloys and a ZrSnMoNb alloy was measured within a temperature interval of 623–823 K, and a broad, applied stress interval. The stress sensitivity parameter, m′ = (∂ ln dot ϵ s /∂ ln σ) T , where σ and T are applied stress and temperature, respectively, was found to increase with applied stress from about 3 at dot ϵ s ⋍ 10 −9 s −1 to eventually more than 20 at dot ϵ s ∼- 10 −5 s −1 . The apparent activation energy of creep, Q, increases with temperature from values lower to values considerably higher than the expected value of activation enthalpy of lattice diffusion, ΔH . Correction of Q for the temperature dependence of the elastic modulus does not lead to the value of the activation energy of creep which is temperature independent and close to ΔH . Nevertheless, it is suggested that creep in the alloys investigated is controlled by recovery at medium applied stresses, while at low applied stresses the modified Nabarro-Herring or Coble mechanism is probably the rate controlling one. At high applied stresses (and lower temperatures) the measured steady state creep rates are significantly contributed to by an athermal deformation mechanism, at least at the lowest tin concentration studied, i.e., 3 wt.%. Creep in the alloys investigated cannot be interpreted in terms of measured effective stress. The influence of alloy composition on steady state creep rate and steady state flow stress, respectively, was evaluated and the creep strengthening effect of tin, molybdenum and niobium is discussed.

Some stress change experiments on the creep of class I alloys

Abstract Some effects of changes in applied stress on the creep of Class I solid solution alloys (Sherby-Burke classification) represented by the AL-5.5 at.% Mg solid solution were investigated. The experimental results obtained are in good agreement with the accepted dislocation model and the concept of effective and internal stress. These results suggest that the moving dislocation density depends on both the applied and the measured effective stress more strongly than on the total dislocation density as determined experimentally.

Some basic creep characteristics of ZrSnMo and ZrSnMoNb alloys part II. Fracture in creep

Abstract In Part II of the present work, some characteristics of fracture in creep of five ZrSnMo alloys and a ZrSnMoNb alloy at temperatures between 673 and 823 K, and in a broad applied stress interval are summarized and discussed. Applied stress and temperature dependences of time to fracture were obtained, as well as relations between creep strain to fracture and time to fracture, and relations between creep strain to fracture and applied stress. Further, relations between “mean” creep rate, defined as the ratio of creep strain to fracture and time to fracture, and steady state creep rate were determined. The results are discussed considering temperature and applied stress dependences of steady state creep rate. It is suggested that fracture in creep is controlled by the same mechanism as the creep itself. The influence of alloy composition on time to fracture was evaluated and the strengthening effect of tin, molybdenum and niobium is discussed.