The Inhibitory Effect of Salmon Calcitonin on Tri-Iodothyronine Induction of Early Hypertrophy in Articular Cartilage

Salmon calcitonin has chondroprotective effect both in vitro and in vivo, and is therefore being tested as a candidate drug for cartilage degenerative diseases. Recent studies have indicated that different chondrocyte phenotypes may express the calcitonin receptor (CTR) differentially. We tested for the presence of the CTR in chondrocytes from tri-iodothyronin (T3)-induced bovine articular cartilage explants. Moreover, investigated the effects of human and salmon calcitonin on the explants.Early chondrocyte hypertrophy was induced in bovine articular cartilage explants by stimulation over four days with 20 ng/mL T3. The degree of hypertrophy was investigated by molecular markers of hypertrophy (ALP, IHH, COLX and MMP13), by biochemical markers of cartilage turnover (C2M, P2NP and AGNxII) and histology. The expression of the CTR was detected by qPCR and immunohistochemistry. T3-induced explants were treated with salmon or human calcitonin. Calcitonin down-stream signaling was measured by levels of cAMP, and by the molecular markers.Compared with untreated control explants, T3 induction increased expression of the hypertrophic markers (p<0.05), of cartilage turnover (p<0.05), and of CTR (p<0.01). Salmon, but not human, calcitonin induced cAMP release (p<0.001). Salmon calcitonin also inhibited expression of markers of hypertrophy and cartilage turnover (p<0.05).T3 induced early hypertrophy of chondrocytes, which showed an elevated expression of the CTR and was thus a target for salmon calcitonin. Molecular marker levels indicated salmon, but not human, calcitonin protected the cartilage from hypertrophy. These results confirm that salmon calcitonin is able to modulate the CTR and thus have chondroprotective effects

Room temperature precipitation in quenched Al-Cu-Mg alloys: a model for the reaction kinetics and yield strength development

The microstructural evolution during low temperature ageing of two commercial purity alloys (Al-1.2Cu-1.2Mg-0.2Mn and Al-1.9Cu-1.6Mg-0.2Mn at.%) was investigated. The initial stage of hardening in these alloys is very rapid, with the alloys nearly doubling in hardness during 20 h ageing at room temperature. The microstructural evolution during this stage of hardening was investigated using differential scanning calorimetry (DSC), isothermal calorimetry and three-dimensional atom probe analysis (3DAP). It is found that during the hardening a substantial exothermic heat evolution occurs and that the only microstructural change involves the formation of Cu-Mg co-clusters. The kinetics of cluster formation is analysed and the magnitude of the hardening is discussed on the basis of a model incorporating solid solution hardening and modulus hardening originating from the difference in modulus between Al and clusters

Precipitates and intermetallic phases in precipitation hardening Al–Cu–Mg–(Li) based alloys

The present study contains a critical review of work on the formation of precipitates and intermetallic phases in dilute precipitation hardening Al–Cu–Mg based alloys with and without Li additions. Although many suggestions for the existence of pre-precipitates in Al–Cu–Mg alloys with a Cu/Mg atomic ratio close to 1 have been made, a critical review reveals that evidence exists for only two truly distinct ones. The precipitation sequence is best represented as: supersaturated solid solution-&gt;co-clusters-&gt;GPB2/S"-&gt;S where clusters are predominantly Cu–Mg co-clusters (also termed GPB or GPB I zones), GPB2/S" is an orthorhombic phase that is coherent with the matrix (probable composition Al10Cu3Mg3) for which both the term GPB2 and S" have been used, and S phase is the equilibrium Al2CuMg phase. GPB2/S" can co-exist with S phase before the completion of the formation of S phase. It is further mostly accepted that the crystal structure of S’ (Al2CuMg) is identical to the equilibrium S phase (Al2CuMg). The Perlitz and Westgren model for S phase is viewed to be the most accepted structure. 3DAP analysis showed that Cu–Mg clusters form within a short time of natural and artificial aging. Cu–Mg clusters and S phase contribute to the first and second stage hardening during aging. In Al–Cu alloys, the theta phase (Al2Cu) has I4/mcm structure with a=0.607 nm and c=0.487 nm, and theta’ phase with tetragonal structure and a=0.404 nm, c=0.58 nm, the space group is I4¯m2. Gerold’s model for theta" (or GPII) appears to be favourable in terms of free energy, and is consistent with most experimental data. The transformation from GPI to GPII (or theta") seems continuous, and as Cu atoms will not tend to cluster together or cluster with vacancies, the precipitation sequence can thus be captured as: supersaturated solid solution-&gt;theta" (Al3Cu)-&gt; theta’ (Al2Cu)-&gt;theta (Al2Cu). The Omega phase (Al2Cu) has been variously proposed as monoclinic, orthorhombic, hexagonal and tetragonal distorted theta phase structures. It has been shown that Omega phase forms initially on {111}Al with c=0.935 nm and on further aging, the c lattice parameter changes continuously to 0.848 nm, to become identical to the orthorhombic structure proposed by Knowles and Stobbs (a=0.496 nm, b=0.858 nm and c=0.848 nm). Other models are either wrong (for example, monoclinic and hexagonal) or refer to a transition phase (for example, the Garg and Howe model with c=0.858 in a converted orthorhombic structure). For Al–Li–Cu–Mg alloys, the L12 ordered metastable delta’ (Al3Li) phase has been observed by many researchers. The Huang and Ardell model for T1 phase (space group P6/mmm, a=0.496 nm and c=0.935 nm), appears more likely than other proposed structures. Other proposed structures are perhaps due to the T1 phase forming by the dissociation of 1/2&lt;110&gt; dislocations into 1/6&lt;211&gt; Shockley partials bounding a region of intrinsic stacking fault, in which copper and lithium enrichment of the fault produces a thin layer of the T1 phase